Magnetic Material and Manufacturing Method Therefor

ABSTRACT

Provided is a new magnetic material with high magnetic stability, as well as a manufacturing method therefor, said magnetic material having a higher saturation magnetization than ferrite-based magnetic materials, and having a higher electrical resistivity than existing metal-based magnetic materials, thus solving problems such as that of eddy current loss. Mn-ferrite nanoparticles obtained through wet synthesis are reduced within hydrogen, and grains are allowed to grow while simultaneously using a phase separation phenomenon due to a disproportionation reaction to produce a magnetic material powder in which an α-(Fe, Mn) phase and a Mn-enriched phase are nano-dispersed. This powder is then sintered to produce a solid magnetic material.

TECHNICAL FIELD

The present invention relates to a soft magnetic material or a semi-hardmagnetic material, and a production method thereof.

BACKGROUND ART

Global environmental problems, such as global warming and exhaustion ofresources, are becoming more severe, and the social demands for energysaving and using less resources in various electronic and electricdevices are increasing day by day. In such a situation, there is a needfor further improvement in the performance of soft magnetic materialsused in the drive unit of motors and the like and the transformer ofvoltage-conversion devices. In addition, to solve various problemsinvolved with manufacturing various compact and high-performanceinformation communication devices, increasing calculation processingspeeds, increasing recording storage capacity, as well as maintainingenvironmental sanitation in infrastructure, distribution systems thatare becoming ever more complex, and strengthening the security thatbecomes increasingly diverse, there is a need to improve theelectromagnetic properties, reliability, and sensitivity of various softmagnetic materials and semi-hard magnetic materials used for variouselements, sensors, and systems.

Demand for next-generation automobiles equipped with large motors drivenat high revolutions (hereinafter, this refers to revolution speedsexceeding 400 rpm) such as in electric automobiles, fuel cellautomobiles, and hybrid automobiles, is expected to further increase inthe future to meet the current calls to deal with environmental andenergy problems. Among the various problems to be solved, betterperformance and lower costs for the soft magnetic material used for thestator in a motor are one of the important issues.

Existing soft magnetic materials used for these applications are broadlydivided into two types, namely, metallic magnetic materials andoxide-based magnetic materials.

Examples of the former, namely, metallic magnetic materials, includesilicon steel (Fe—Si), which is a Si-containing crystalline materialbeing a typical example of electromagnetic steels, as well as sendust(Fe—Al—Si), which is an intermetallic compound containing Al,electromagnetic soft iron (Fe), which is pure iron having a low carboncontent of 0.3% by mass or less and a low impurity content, amorphousalloys such as permalloy, which contains Fe—Ni as a main component, andMetglas (Fe—Si—B), and a group of nanocrystalline soft magneticmaterials (whose representative compositions include Fe—Cu—Nb—Si—B,Fe—Si—B—P—Cu, etc.), such as Finemet, which are nanocrystal-amorphousphase-separated materials obtained by precipitating microcrystals byapplying an appropriate heat treatment to the amorphous alloy. The term“nano” as used here means a size of 1 nm or more and less than 1 μm. Formagnetic materials other than nanocrystalline soft magnetic materials,in terms of reducing coercive force and iron loss, it is important tofacilitate movement of the domain walls as a composition that is asuniform as possible. It is noted that nanocrystalline soft magneticmaterials are a heterogeneous system that include a crystalline phase,an amorphous phase, a Cu-enriched phase, and the like, and magnetizationreversal is considered to be mainly caused by magnetization rotation.

Examples of the latter, namely, oxide-based magnetic materials, includeferritic magnetic materials such as Mn—Zn ferrite and Ni—Zn ferrite.

Silicon steel has until now been the soft magnetic material that is mostwidely used in high-performance soft magnetic material applications, andis a high magnetization, low coercive force magnetic material having asaturation magnetization of 1.6 to 2.0 T and a coercive force of 3 to130 A/m. This material is obtained by adding up to 4% by mass of Si toFe, which lowers the magnetocrystalline anisotropy and the saturationmagnetostriction constant and reduces the coercive force withoutsignificantly impairing the large magnetization of Fe. In order toimprove the performance of this material, it is necessary to removeforeign substances that hinder the movement of domain walls whileincreasing the crystal grain size by appropriately combiningcomposition-controlled materials with the appropriate hot and coldrolling and annealing. In addition to non-oriented steel sheets with arandom orientation of the crystal grains, directional steel sheets inwhich the (100) direction of Fe—Si, which is an easily magnetizeddirection, is highly oriented in the rolling direction are widely usedas a material that further reduces coercive force.

Since this material is a rolled material, it has a thickness of lessthan about 0.5 mm. Further, since this material is a homogeneous metalmaterial, it has a low electric resistance of about 0.5 μΩm. Generally,this material is used in large equipment applications by covering thesurface of each silicon steel sheet with an insulating film, punchingout with a die, and laminating and welding to provide thickness whilesuppressing eddy current loss that occurs in high-rotation applications,such as next-generation automobiles. Therefore, the costs of thepunching and lamination steps, and deterioration of the magneticproperties are serious problems.

A nanocrystalline soft magnetic material such as Fe—Cu—Nb—Si—B is a softmagnetic material having a nanocrystalline structure in which theamorphous grain boundary phases are randomly oriented that is obtainedby subjecting an alloy which has become amorphous by rapid cooling to aheat treatment at a temperature higher than the crystallizationtemperature to cause crystal grains of about 10 nm to precipitate in theamorphous phase. The coercive force of this material is extremely low,namely, 0.6 to 6 A/m, and the saturation magnetization is 1.2 to 1.7 T,which is higher than that of an amorphous material. Hence, the marketfor such materials is expanding at present. This material is arelatively new material that was developed in 1988. The principle behindthese magnetic properties is that by making the crystal grain sizesmaller than the ferromagnetic exchange length (also called the exchangecoupling length) and by causing the randomly-oriented main phase,namely, the ferromagnetic phase, to undergo ferromagnetic couplingthrough an amorphous interface phase, the magnetocrystalline anisotropyis averaged, thereby reducing the coercive force. This mechanism iscalled a random magnetic anisotropy model, or a random anisotropy model(e.g., see Non Patent Document 1).

However, since this material is produced after first producing a ribbonby liquid rapid quenching, the thickness of the product is about 0.02 to0.025 mm, and the insulation, cutting, alignment, lamination, welding,and annealing steps are more complicated than for silicon steel, thismaterial suffers from problems such as processability and thermalstability. Furthermore, the electric resistivity is small at 1.2 μΩm,and a problem with eddy current loss similar to other rolled materialsand ribbons has been pointed out.

In order to overcome this, attempts have been made to prepare a bulkmolding material by pulverizing the above-mentioned ribbon-shapednanocrystalline soft magnetic material using SPS (discharge plasmasintering) (e.g., see Non Patent Document 2). However, the magneticproperties are much worse than for a 20 μm ribbon, with a coercive forceof 300 A/m and a saturation magnetization of 1 T. At present, there isno good method other than a lamination method for producing productsthicker than 0.5 mm.

Among existing soft magnetic materials so far, ferrite oxide materialshave the least problems with eddy current loss in high-rotationapplications. The electric resistivity of such a material is 10⁶ to 10¹²μΩm, and the material can be easily bulked to 0.5 mm or more bysintering. Further, such a material can also be formed into a moldedbody free from eddy current loss. Therefore, it is a material suitablefor high-rotation, high-frequency applications. In addition, since it isan oxide, this material does not rust and the stability of its magneticproperties is also excellent. However, the coercive force of thismaterial is comparatively high, namely, 2 to 160 A/m, and in particular,the saturation magnetization is small at 0.3 to 0.5 T. Therefore, thismaterial is not suitable for high-performance, high-speed motors fornext-generation automobiles, for example.

In general, metallic soft magnetic materials such as silicon steel areused by laminating due to their thin thickness as a result of beingrolled materials. However, such materials have a low electricresistance, and suffer from the occurrence of eddy current loss forhigh-rotation, high-performance motors. Consequently, lamination needsto be carried out in order to solve these problems. This results inserious problems such as the steps becoming complicated, an insulationtreatment before lamination being necessary, magnetic propertiesdeteriorating due to punching and the like, and increased costs for thesteps. On the other hand, oxide-based soft magnetic materials such asferrite have a large electric resistance and no problems with eddycurrent loss, but they are unsuitable for high-performance motors fornext-generation automobiles because they have a small saturationmagnetization of 0.5 T or less. However, from the perspective ofoxidation resistance, oxide-based soft magnetic materials are superiorto metallic soft magnetic materials in terms of having a high stability.

The upper limit of the thickness that can be used for the motor in thenon-oriented electromagnetic steel sheets of silicon steel that areproduced for high-performance motors for next-generation automobilesusing permanent magnets is estimated as follows. First, when analternating magnetic field with a frequency f is applied to thematerial, a skin depth s at which the strength of the magnetic field is1/e is as shown in the following relational expression (1).

$\begin{matrix}\left\lbrack {{Expression}\mspace{14mu} 1} \right\rbrack & \; \\{s = \sqrt{\frac{\rho}{\pi \; {\mu\mu}_{0}f}}} & (1)\end{matrix}$

In the case of a non-oriented electromagnetic steel sheet, which issilicon steel, by substituting an electric resistivity p=5×10⁻⁷ [Ωm] anda permeability μμ₀=9200×4π×10⁻⁷ [N/A²] as representative values intorelational expression (1), when the number of poles of thenext-generation automotive motor is 8 and the maximum rotation speed is10000 rpm, namely, f is 667 [Hz], the skin depth is 0.14 mm.

The condition for preventing a large reduction in the effectivemagnetization of the material is to set the particle size of thematerial to no greater than twice the skin depth. Therefore, forexample, when a silicon steel sheet is used at 667 Hz, the sheetthickness needs to be about 0.3 mm, but since the thickness of thenext-generation automotive motor is, for example, 9 cm, when a thinsilicon steel sheet having a thickness of 0.3 mm is used, about 300sheets each have to be insulated and laminated. The steps of insulating,punching, aligning, welding and annealing such a thin sheet arecomplicated and expensive. In order to make the laminated sheetthickness as thick as possible, it is necessary to increase the electricresistivity of the material.

Thus, there is a need for the appearance of a soft magnetic materialhaving an electric resistance higher than a metallic silicon steel sheetand the like, a saturation magnetization higher than a ferrite magneticmaterial, and physical properties to compensate for problems caused byhaving such an electric resistance and saturation magnetization, namely,a soft magnetic material combining both the advantages of the highsaturation magnetization of metallic magnetic materials and the smalleddy current loss, absence of a lamination step and complicated stepsassociated therewith, high oxidation resistance, and good magneticstability of oxide-based magnetic materials.

PRIOR ART DOCUMENTS Non Patent Document [Non Patent Document 1]

G. Herzer, IEEE Transactions on Magnetics, vol. 26, No. 5 (1990) pp.1397-1402

[Non Patent Document 2]

Y Zhang, P. Sharma and A. Makino, AIP Advances, vol. 3, No. 6(2013)062118

SUMMARY OF INVENTION Technical Problem

It is an object of the present invention to provide, by using a magneticmaterial in which an α-(Fe,Mn) phase and a Mn-enriched phase arenano-dispersed, a new magnetic material, and a production methodthereof, having high magnetic stability, which enables a highersaturation magnetization to be realized than a ferrite magneticmaterial, and enables the above-mentioned problems such as eddy currentloss to be solved due to having a higher electric resistivity thanexisting metallic magnetic materials.

Further, it is an object of the present invention to provide a powdersintered magnetic material that is capable of producing a molded bodyhaving a thickness of 0.5 mm or more, further 1 mm or more, and even 5mm or more, by simple steps without performing complicated steps such aslamination, and at the same time can reduce eddy current.

Solution to Problem

The present inventors extensively studied magnetic materials thatsimultaneously satisfy two points, namely, having a high magnetizationand being able to solve the above-mentioned problem of eddy current lossdue to a high electric resistivity, which are contradictorycharacteristics for conventional magnetic materials, and yet haveexcellent electromagnetic properties that combine the merits of bothmetallic magnetic materials and oxide-based magnetic materials which donot require complicated steps such as lamination, as well as have stablemagnetic properties even in air. As a result, the present inventorsdiscovered that a magnetic material containing two or more of variouscrystal phases, or one kind of crystal phase and an amorphous phase, canbe obtained through disproportionation during a reduction reaction ofmanganese ferrite (in the present invention, also referred to as“Mn-ferrite”), which is completely different from theconventionally-used uniform homogeneous crystalline and amorphousmaterials or, among amorphous materials, nanocrystalline soft magneticmaterials in which uniform nanocrystals are precipitated, and completedthe present invention by controlling the composition, the crystalstructure, the crystal grain size, and the powder particle diameter ofthe magnetic material, establishing a method for producing the magneticmaterial, and establishing a method for solidifying the magneticmaterial without laminating.

In order to solve the above problem, there is a need for a magneticmaterial having a saturation magnetization that is 0.3 T, namely, sincethe density of the magnetic material of the present invention is closeto the density of a metal system, the saturation magnetization needs tobe at a level of 30 emu/g or higher when calculated in terms of thedensity of Fe, and an electric resistivity of 1.5 μΩm or more. Inparticular, just in terms of a soft magnetic material, the saturationmagnetization needs to be preferably 100 emu/g or more, and morepreferably 150 emu/g or more.

Specifically, the present invention is as follows.

(1) A soft magnetic or semi-hard magnetic material, comprising a firstphase having crystals with a bcc structure containing Fe and Mn and asecond phase containing Mn, the second phase having a Mn content thatis, when a sum of the Fe and the Mn contained in the second phase istaken to be 100 atom %, larger than a Mn content when a sum of the Feand the Mn contained in the first phase is taken to be 100 atom %.(2) The magnetic material according to the above (1), wherein themagnetic material is soft magnetic.(3) The magnetic material according to the above (1) or (2), wherein thefirst phase has a composition represented by a composition formulaFe_(100-x)Mn_(x) (where x is 0.001≤x≤33 in terms of atomic percentage).(4) The magnetic material according to any one of the above (1) to (3),wherein the first phase has a composition represented by a compositionformula Fe_(100-x)(Mn_(100-y)M_(y))_(x/100) (where x and y are0.001≤x≤33 and 0.001≤y≤50 in terms of atomic percentage, and M is one ormore of Zr, Hf, Ti, V, Nb, Ta, Cr, Mo, W, Ni, Co, Cu, Zn, and Si).(5) The magnetic material according to any one of the above (1) to (4),wherein the second phase is a phase having crystals with a bcc structurecontaining Fe and Mn, and a Mn content, when the sum of the Fe and theMn contained in the phase is taken to be 100 atom %, is an amount of 2times or more and 10⁵ times or less and/or 2 atom % or more and 100 atom% or less relative to the Mn content when the sum of the Fe and the Mncontained in the first phase is taken to be 100 atom %.(6) The magnetic material according to any one of the above (1) to (5),wherein the second phase comprises a Mn-ferrite phase.(7) The magnetic material according to any one of the above (1) to (6),wherein the second phase comprises a wustite phase.(8) The magnetic material according to any one of the above (1) to (7),wherein the phase having crystals with a bcc structure containing Fe andMn has a volume fraction of 5% by volume or more based on the wholemagnetic material.(9) The magnetic material according to the above (6) or (7), wherein themagnetic material has a composition in a range of, based on thecomposition of the whole magnetic material, 20 atom % or more and 99.998atom % or less of Fe, 0.001 atom % or more and 50 atom % or less of Mn,and 0.001 atom % or more and 55 atom % or less of O.(10) The magnetic material according to any one of the above (1) to (9),wherein an average crystal grain size of the first phase, the secondphase, or the whole magnetic material is 1 nm or more and 10 μm or less.(11) The magnetic material according to any one of the above (1) to(10), wherein at least the first phase has a bcc phase having acomposition represented by a composition formula Fe_(100-x)Mn_(x) (wherex is 0.001≤x≤1 in terms of atomic percentage), and that bcc phase has acrystallite size of 1 nm or more and less than 100 nm.(12) The magnetic material according to any one of the above (1) to(11), wherein the magnetic material is in a powder form, and when themagnetic material is soft magnetic, the magnetic material has an averagepowder particle diameter of 10 nm or more and 5 mm or less, and when themagnetic material is semi-hard magnetic, the magnetic material has anaverage powder particle diameter of 10 nm or more and 10 μm or less.(13) The magnetic material according to any one of the above (1) to(12), wherein at least the first phase and the second phase areferromagnetically coupled with adjacent phases.(14) The magnetic material according to any one of the above (1) to(13), wherein the first phase and the second phase are continuouslybonded to each other directly or via a metal phase or an inorganic phaseto form a massive state as the whole magnetic material.(15) A method for producing the magnetic material according to the above(12) by reducing a manganese ferrite powder having an average powderparticle diameter of 1 nm or more and less than 1 μm in a reducing gascontaining hydrogen gas at a reduction temperature of 400° C. or moreand 1350° C. or less.(16) A method for producing the magnetic material according to any oneof the above (1) to (13) by reducing a manganese ferrite powder havingan average powder particle diameter of 1 nm or more and less than 1 μmin a reducing gas containing hydrogen gas, and forming the first phaseand the second phase by a disproportionation reaction.(17) A method for producing the magnetic material according to the above(14) by sintering the magnetic material produced by the method accordingto the above (15) or (16).(18) A method for producing a soft magnetic or semi-hard magneticmaterial, comprising performing annealing at least once after thereduction step in the method according to the above (15), or after thereduction step or the formation step in the method according to theabove (16), or after the sintering step in the method according to theabove (17).

Advantageous Effects of Invention

According to the present invention, there can be provided a magneticmaterial having a high saturation magnetization and a small eddy currentloss, in particular a soft magnetic material that is suitably used evenin high rotation motors and the like, and various soft magneticmaterials and semi-hard magnetic materials having high oxidationresistance.

According to the present invention, because the magnetic material can beused in the form of a powder material like ferrite, it can easily beproduced in bulk by sintering or the like, and hence the presentinvention can solve problems such as complicated steps like laminationand the like caused by the use of metallic soft magnetic materials knownas thin sheets, as well as the high costs involved with such steps.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is an SEM image of a powder (Example 13) obtained by reducing an(Fe_(0.994)Mn_(0.006))₄₃O₅₇ ferrite nanopowder in hydrogen gas at 900°C. for 1 hour.

FIG. 2(A) and FIG. 2(B) are both an SEM image of a powder (Example 5)obtained by reducing an (Fe_(0.672)Mn_(0.328))₄₃O₅₇ ferrite nanopowderin hydrogen at 650° C. for 1 hour (the numerical values in the diagramsare the Mn content of each phase).

FIG. 3 is an SEM image of a magnetic material powder (Example 10)obtained by reducing an (Fe_(0.672)Mn_(0.328))₄₃O₅₇ ferrite nanopowderin hydrogen at 1100° C.

FIG. 4 is an SEM image of an (Fe_(0.672)Mn_(0.328))₄₃O₅₇ Mn-ferritenanopowder (Comparative Example 1).

FIG. 5 is an SEM image of a magnetic material powder (Example 1)obtained by reducing an (Fe_(0.672)Mn_(0.328))₄₃O₅₇ ferrite nanopowderin hydrogen at 900° C. (the numerical values in the diagrams are the Mncontent at the +position).

FIG. 6 is an X-ray diffraction diagram of an α-Fe powder (ComparativeExample 4) and powders produced by reducing a Mn-ferrite nanopowderhaving an (Fe_(0.672)Mn_(0.328))₄₃O₅₇ composition in hydrogen at varioustemperatures of 600° C. or more and up to 1000° C. (Example 4 (reductiontemperature=600° C.), Example 7 (reduction temperature=800° C.), andExample 9 (reduction temperature=1000° C.)).

FIG. 7 shows a reduction temperature dependence of saturationmagnetization (emu/g) and coercive force (kA/m) for an Fe—Mn magneticmaterial powder (Examples 1 to 11 and Comparative Example 5).

FIG. 8 is an SEM image of an (Fe_(0.994)Mn_(0.006))₄₃O₅₇ Mn-ferritenanopowder (Comparative Example 6).

FIG. 9 is an SEM image of the surface of the Fe_(71.1)Mn_(28.9) solidmagnetic material of Example 19.

FIG. 10 is an oxygen characteristic X-ray surface distribution map ofthe surface of the Fe_(71.1)Mn_(28.9) solid magnetic material of Example19.

FIG. 11 is a Mn characteristic X-ray surface distribution map of across-section of the Fe_(99.9)Mn_(0.1) powder of Example 20.

FIG. 12 is a histogram representing a distribution of the Mn content ofa cross-section of the Fe_(99.9)Mn_(0.1) powder of Example 20.

FIG. 13 is a Mn characteristic X-ray surface distribution map of across-section of the Fe_(70.2)Mn_(29.8) powder of Example 21.

FIG. 14(A) and FIG. 14(B) are diagrams plotting a correlation between Mncontent and Fe or O content at each EDX measurement point on across-section of the Fe_(70.2)Mn_(29.8) powder of Example 21 (FIG.14(A): diagram plotting a correlation between Mn content and Fe content,and FIG. 14(B): diagram plotting a correlation between Mn content and Ocontent).

DESCRIPTION OF EMBODIMENTS

The present invention will now be described in detail.

The term “magnetic material” as used in the present invention refers tomagnetic materials called “soft magnetic” (i.e., “soft magneticmaterials”) and magnetic materials called “semi-hard magnetic” (i.e.,“semi-hard magnetic materials”). Here, a “soft magnetic material” asreferred to in the present invention is a magnetic material having acoercive force of 800 A/m (≈10 Oe) or less, and a “semi-hard magneticmaterial” as referred to in the present invention is a magnetic materialhaving a coercive force exceeding 800 A/m and 40 kA/m (≈500 Oe) or less.In order to obtain an excellent soft magnetic material, it is importantto have a low coercive force, a high saturation magnetization orpermeability, and low iron loss. The causes of iron loss are mainlyhysteresis loss and eddy current loss. In order to reduce the former, itis necessary to make the coercive force smaller, and in order to reducethe latter it is necessary for the electric resistivity of the materialitself to be high, or to increase the electric resistance of the wholemolded body to be subjected to practical use. For a semi-hard magneticmaterial, it is required to have a coercive force that is appropriatefor the application, and to have a high saturation magnetization andresidual magnetic flux density. Among magnetic materials, soft magneticor semi-hard magnetic materials used for high frequency generate a largeeddy current, and hence it is important for the material to have a highelectrical resistivity and that the powder particle diameter is small,or the sheet thickness is a thin sheet or ribbon.

The term “ferromagnetic coupling” as used in the present inventionrefers to a state in which adjacent spins in a magnetic material arestrongly bound by exchange interaction. In particular in the presentinvention, this term refers to state in which the spins of two adjacentcrystal grains (and/or amorphous grains) are strongly bound to eachother by exchange interaction across the crystal boundary. Sinceexchange interaction is an interaction that only reaches a distancebased on the short range order of the material, when a nonmagnetic phaseis present at the crystal boundary, exchange interaction does not workon the spins in the region on either side thereof, and henceferromagnetic coupling does not occur between the crystal grains (and/oramorphous grains) on either side. In the present application, the term“crystal grain” may include amorphous grains. Further, thecharacteristics of the magnetic curve of the material in whichferromagnetic coupling has occurred between different adjacent crystalgrains having different magnetic properties will be described later.

The term “disproportionation” as used in the present invention meansthat phases having two or more different compositions or differentcrystal structures are produced from a phase in a homogeneouscomposition by a chemical reaction. In the present invention,disproportionation is caused as a result of a reducing substance such ashydrogen being involved in a phase of the homogeneous compositionleading to the occurrence of a reduction reaction. During this“disproportionation” reaction, water is often produced as a by-product.

In the present invention, the expression “including an Fe component anda Mn component” means that the magnetic material of the presentinvention always contains Fe and Mn as components, and optionally the Mnmay be substituted with a certain amount of other atoms (specifically,one or more of Zr, Hf, Mn, V, Nb, Ta, Cr, Mo, W, Ni, Co, Cu, Zn, andSi). Further, oxygen (O component) may be contained, and when an Ocomponent or iron oxide hydroxide, or the like is present as a minorphase, H may be contained mainly as an OH group, and other unavoidableimpurities as well as Cl or alkali metals such as K derived from rawmaterials may also be included. Alkali metals such as K are suitablecomponents in that they may exert an effect of promoting the reductionreaction.

The term “magnetic powder” generally refers to a powder havingmagnetism, but in the present application a powder of the magneticmaterial of the present invention is referred to as “magnetic materialpowder”. Therefore, the term “magnetic material powder” is included inthe term “magnetic powder”.

The present invention relates to a magnetic material comprising a phase(first phase) containing manganese in an α-Fe phase and a Mn-enrichedphase (second phase) having a Mn content higher than the first phase.The best mode of the present invention is a “powder” in which the twophases are mixed and bonded at the nano level. These magnetic materialpowders are used for various devices by directly compacting orsintering. Further, depending on the application, an organic compoundsuch as a resin, an inorganic compound such as glass or ceramic, acomposite material thereof, or the like may be added and the resultantmixture may be molded.

Hereinafter, the composition, crystal structure and morphology, crystalgrain size and powder particle diameter, and the production method ofthe first phase containing Fe and Mn and the second phase enriched withMn will be described. In particular, a method for producing ananocomposite oxide powder as a precursor of the magnetic material ofthe present invention, a method for reducing the powder, a method forsolidifying the reduced powder, and a method for annealing in each stepof these manufacturing methods, will be described.

<First Phase>

In the present invention, the first phase is a crystal having a bccstructure cubic crystal (space group Im-3m) containing Fe and Mn as acrystal structure. The Mn content of this phase is preferably 0.001 atom% or more and 33 atom % or less, based on a sum (total content) of theFe and Mn contained in the phase of 100 atom %. Specifically, thecomposition of the first phase may be represented by the compositionformula Fe_(100-x)Mn_(x) (where x is 0.001≤x≤33 in terms of atomicpercentage).

Here, the Mn content and the Fe content are, unless stated otherwise,respectively the value of the atomic ratio of Mn or Fe relative to thesum (in the present application, as described above, sometimes alsoreferred to as the total content, or as total amount) of Fe and Mncontained in the phase. In the present invention, this may berepresented as an atomic percentage relative to the sum (total content)of Fe and Mn contained in the phase of 100 atom %.

It is preferable that the Mn content is 33 atom % or less in order tosuppress a decrease in magnetization. Further, the Mn content is morepreferably 20 atom % or less, because this means that, depending on theproduction method and conditions, a magnetization exceeding 1 T can berealized. In addition, the Mn content is even more preferably 10 atom %or less, as this enables a magnetic material having a saturationmagnetization exceeding 1.6 T to be produced. Further, the Mn content ispreferably 0.001 atom % or more, as this means that, unlike when Fe isused alone, the magnetic properties in the soft magnetic region can beadjusted by the effect of Mn addition. A particularly preferable rangefor the Mn content is 0.01 atom % or more and 10 atom % or less. In thisregion, depending on the production conditions, it is possible toprepare a semi-hard magnetic material from soft magnetic, and themagnetic material has even more preferable electromagnetic properties.Even if the coercive force is slightly sacrificed, when it is desired toproduce a soft magnetic material having a higher magnetization, it ispreferable to set the Mn content of the first phase to 5 atom % or less.

The first phase having a Fe—Mn composition having this bcc structure isalso referred to as an α-(Fe,Mn) phase in the present application, sincethe symmetry of the crystal is the same as the α-phase, which is theroom temperature phase of Fe.

When the content of the Mn component of the first phase of the presentinvention is taken to be 100 atom %, 0.001 atom % or more and less than50 atom % of the Mn can be with one or more of any of Zr, Hf, Ti, V, Nb,Ta, Cr, Mo, W, Ni, Co, Cu, Zn, and Si (in the present application, thesesubstitution elements are also referred to as “M component”). Therefore,in the present invention, when the Mn contained in the first phase has acomposition substituted with the M component, the combination of the Mnand the M component in the composition is equivalent to theabove-mentioned “Mn component”, and the Mn component content(specifically, the sum of the Mn content and the M component content inthe composition) is 100 atom %. Among these M components, co-adding alarge number of elemental species to the soft magnetic material of thepresent invention is effective in reducing coercive force. Inparticular, in terms of atomic percentage when the Mn component contentof the first phase is taken to be 100 atom %, containing 1 atom % ormore of Ti, V, Cr, and Mo is effective in enabling the nanocrystals ofthe present invention to be easily produced without largely depending onthe cooling rate in the reduction treatment and the annealing treatment.Further, since Zr, Hf, Ti, Cr, V, Ni, Co, and Si decrease theanisotropic magnetic field, they are preferable as components coexistingwith the soft magnetic material of the present invention. Also, in orderto improve the saturation magnetization, Ni is preferably added in anamount of about 5 atom % or less, and Co is preferably added in anamount of less than about 50 atom %. Zr, Hf, Ti, V, Nb, Ta, Cr, Mo, andW suppress improper grain growth during the reduction step even when 1atom % or less is added in terms of atomic percentage when the Mncomponent content of the first phase is taken to be 100 atom %. Ti, Ni,Co, Cu, and Zn are preferable for improving oxidation resistance andmolding properties. Of these, when Ti is co-added to Mn, not only theabove effects but also a unique synergistic effect is realized in whicha low coercive force and high magnetization are achieved. The preferablesubstitution quantity of Ti for Mn is 0.01 atom % or more and less than50 atom %. A more preferable M component content is not dependent on theelemental species, and is 0.1 atom % or more and 30 atom % or less interms of the substitution quantity for Mn.

Therefore, for example, when the first phase has a composition formulaFe_(100-x)Mn_(x) (where x is 0.001≤x≤33 in terms of atomic percentage),if its Mn component is substituted with the M component in the range of0.01 atom % or more and less than 50 atom %, the composition formulathereof is represented as Fe_(100-x)(Mn_(100-y)M_(y))_(x/100) (where xand y are 0.001≤x≤33 and 0.001≤y≤50 in terms of atomic percentage, and Mis one or more of Zr, Hf, Ti, V, Nb, Ta, Cr, Mo, W, Ni, Co, Cu, Zn, andSi).

A more preferable M component content is not dependent on the elementalspecies, and is 0.1 atom % or more and 30 atom % or less in terms of thesubstitution quantity for Mn.

Note that “improper grain growth” means that the nano-microstructure ofthe magnetic material of the present invention collapses and crystalgrains grow with a homogeneous crystal structure. On the other hand,“grain growth” that is suitable in the present invention is growth inwhich the powder particle diameter grows to be large while maintainingthe nano-microstructure that is a characteristic of the presentinvention, or growth in which a nano-microstructure appears in thecrystal due to a disproportionation reaction, phase separation or thelike after the powder particle diameter has grown to be large, or bothof these cases. Unless otherwise noted, the term “grain growth” in thepresent invention refers to grain growth that is not improper and thatcan generally be said to be suitable. Even when the grain growth isimproper or suitable, the surface area of the magnetic material per unitmass or per unit volume becomes small, and hence oxidation resistancegenerally tends to be improved.

For any of the M components, from the perspective of the addition effectdescribed above, the added amount is preferably 0.001 atom % or more interms of the atomic percentage when the Mn component content of thefirst phase is taken to be 100 atom %, and from the perspective ofpreventing inhibition of the various effects of the Mn component in themagnetic material of the present invention, the added amount ispreferably less than 50 atom %. In the present invention, when expressedas “Mn component”, or when expressed as “Mn”, “manganese” in the contextof discussing formulas such as “α-(Fe,Mn)” phase or the composition ofthe magnetic material, the present invention includes not only cases inwhich Mn is used alone, but also compositions in which 0.001 atom % ormore and less than 50 atom % of the Mn content is substituted with an Mcomponent. In addition, although it is necessary to remove as much aspossible impurities mixed in during the steps, unavoidable impuritiessuch as H, C, Al, Si, S, N, alkali metals such as Li, K and Na, alkaliearth metals such as Mg, Ca, Sr, rare earth metals, halogens such as Cl,F, Br, I, and the like may be included. However, the content of suchimpurities is to be 5 atom % or less, preferably 2 atom % or less, morepreferably 0.1 atom % or less, and particularly preferably 0.001 atom %or less, of the whole (i.e., sum of Fe and Mn contained in the firstphase). This is because the greater the content of these impurities, thelower the magnetization, and in some cases, the coercive force is alsoadversely affected, which depending on the application may deviate fromthe target range. On the other hand, when some components, such asalkali metals like K, which act as reducing aids if contained to someextent, are contained in an amount of 0.0001 atom % or more and 5 atom %or less of the total (i.e., sum of Fe and Mn contained in the firstphase), a magnetic material having a high saturation magnetization maybe obtained. Therefore, when the above-mentioned impurities hinder theobject of the present invention, it is most desirable not to includesuch impurities.

The α-Fe phase not containing Mn is not included in the first phase orthe second phase. The reason for this is that if the content of elementsother than Mn is also extremely small, the α-Fe phase not containing Mnis expected to have saturation magnetization like electromagnetic softiron, but even if the α-Fe phase is a powder in the nano region, theeffect on electric resistivity is poor, oxidation resistance is poor,and the material is inferior in cutting processability. However, theα-Fe phase not containing Mn may exist as a separate phase as long as itdoes not hinder the object of the present invention. When the presentinvention is a soft magnetic material, the volume fraction of the α-Fephase is preferably less than 50% by volume based on the whole magneticmaterial of the present invention.

The volume fraction referred to here is the ratio of the volume occupiedby the target component based on the total volume of the magneticmaterial.

<Second Phase>

In the present invention, the second phase is a phase in which the Mncontent relative to the sum of Fe and Mn contained in the phase islarger than the Mn content relative to the sum of Fe and Mn contained inthe first phase. In other words, in the present invention, the secondphase is a phase in which the atomic percentage of Mn relative to thesum of Fe and Mn contained in the phase is larger than the atomicpercentage of Mn relative to the sum of Fe and Mn contained in the firstphase. Examples of the second phase may include cubic crystals such asan α-(Fe_(1-y)Mn_(y)) phase (phase in space group Im-3m having the samecrystal phase as the first phase, but a higher Mn content than the firstphase), a γ-(Fe,Mn) phase (space group Fm-3m), a wustite phase(representative composition is an (Fe_(1-z)Mn_(z))_(a)O phase, a isusually 0.83 to 1, and is a solid solution of FeO and MnO; in thepresent specification this phase is sometimes simply referred to as an(Fe,Mn)O phase or an (Mn,Fe)O phase; in the present invention, unlessstated otherwise, the term wustite simply refers to a compositioncontaining manganosite, in which 0≤z≤1), a MnO phase (manganositephase), a Mn-ferrite phase (representative composition is(Fe_(1-w)Mn_(w))₃O₄ phase), an α-(Mn_(v)Fe_(1-v)) phase (α-Fe structureis an A2 type in the Strukturbericht classification, with a phasebelonging to the space group of Im-3m, but the phase described here issyngeneous with α-Mn, and the elemental species is an A12 type in theStrukturbericht classification in the case of Mn alone of elementalspecies, belonging to the I4-3m space group (overline above the 4), andis sometimes simply referred to as an α-(Mn,Fe) phase in the presentspecification), a β-(Mn,Fe) phase (space group P4₁32), aγ-(Fe_(1-a)Mn_(a))₂O₃ phase (having a crystal structure different from arhombohedral Mn-hematite phase), a γ-Mn₂O₃ phase, and a Mn₃O₄ phase,cubic crystals such as a Mn(OH)₂ phase, rhombohedral crystals such as anα-(Fe_(1-b)Mn_(b))₂O₃ phase (Mn-hematite phase), an (Mn,Fe)CO₃ phase(hexagonal manganese ore phase, eutectic phase of hexagonal manganeseore and hexagonal siderite), tetragonal crystals such as, among MnO₂, arutile phase, orthorhombic crystals such as a MnO(H) phase (manganitephase, grout ore phase), as well as a Mn—Fe amorphous phase and thelike, and mixtures thereof. It is noted that regarding the Mn—Feamorphous phase, although this depends on the Mn content and reductionconditions, when this phase is present, microcrystals such as theexisting nanocrystal-amorphous phase separation type material describedabove do not form a fine structure being in the form of islands andfloating in an amorphous sea, but often exist in an island shapeseparated from the first phase. The content of the Mn—Fe amorphous phaseis between 0.001 and 10% by volume, and from the perspective ofsuppressing a reduction in magnetization, preferably not more than this.To obtain a magnetic material with higher magnetization, this content ispreferably 5% by volume or less. The amorphous phase and the like may becontained in order to control the disproportionation reaction itself,but in this case, it is preferable to set the content to more than0.001% by volume from the perspective of controlling this reaction.

Although the second phase is in most cases inferior to the first phasein terms of saturation magnetization, when these phases coexist, theelectric resistivity is greatly increased. Further, in the presentinvention, the coercive force is improved when forming the semi-hardmagnetic material. Conversely, in the present invention, when forming asoft magnetic material, depending on the crystal structure, composition,microstructure, interface structure and the like of the phases, it ispossible to realize a small coercive force by ferromagnetically couplingwith the phases. In addition, in the second phase as well, similarly tothe first phase, it is possible to substitute less than 50 atom % of theMn content (wherein the content of the Mn component in the second phaseis taken to be 100 atom %) with an M component. Here, in the presentinvention, similarly to the above-mentioned “Mn component” of the firstphase, the “Mn component” of the second phase also refers to, when theMn contained in the second phase is substituted with an M component, thecombination of the Mn and the M component in its composition.

<Minor Phase, Other Phases>

A phase that does not contain Fe or Mn, and that is mixed with only an Mcomponent compound, is not included in the first phase or the secondphase. However, there are cases where such a phase contributes toimproving electric resistivity, oxidation resistance, sinterability, andthe electromagnetic properties of the semi-hard magnetic material of thepresent invention. In the present application, a phase that does notcontain a Mn component, such as a compound phase of the above-mentionedM component or an Fe compound phase, and a phase in which the content ofthe M component is equal to or more than the content of the Mn element,is referred to as a “minor phase”.

Other than the first phase and the second phase, the magnetic materialmay also contain a minor phase that does not contain Mn, such as awustite phase, a magnetite phase (Fe₃O₄), a maghemite phase (γ-Fe₂O₃), ahematite phase (α-Fe₂O₃), an α-Fe phase, and a γ-Fe phase, an iron oxidehydroxide phase that may or may not contain Mn, such as goethite,akagenite, lepidocrocite, ferroxyhyte, ferrihydrite, green rust, ahydroxide such as potassium hydroxide and sodium hydroxide, a chloridesuch as sodium chloride and potassium chloride, a fluoride, a carbide, anitride, a hydride, a sulfide, a nitrate, a carbonate, a sulfate, asilicate, a phosphate, and the like. The volume of the above minor phaseor the like needs to be smaller than the total volume of the α-(Fe,Mn)phases in the first phase or in the first phase and the second phase inorder for the magnetic material of the present invention to have a highsaturation magnetization and also to exhibit stable magnetic propertiesand high magnetization over time. From the perspective of suppressing adecrease in the saturation magnetization, the preferable range of thecontent of these phases is 50% by volume or less based on the volume ofthe whole magnetic material.

The content of the M component in all of the phases, including the firstphase, the second phase, and the minor phase, must not exceed the Mncontent contained in the first phase and the second phase based on allthe phases. This is because when the content of the M component exceedsthe Mn content, the unique characteristic effects on electromagneticproperties specific to Mn, for example, improved magnetization when asmall amount is added and suppression of a decrease in magnetizationwhen more than that amount is added, improvement in electricresistivity, a remarkable effect on oxidation resistance, and the like,are lost. In the present application, the Mn content of the first phaseand/or the second phase is an amount that includes the M component.

<Case in which Second Phase has Same Crystal Structure as First Phase>

Although the second phase may have the same crystal structure as thefirst phase, it is important that the phases are sufficiently differentfrom each other in terms of their composition. For example, it ispreferable that the Mn content of the second phase relative to the sumof Fe and Mn in the second phase is either twice or more the Mn contentof the first phase relative to the sum of Fe and Mn in the first phase,or the Mn content of the second phase relative to the sum of Fe and Mnin the second phase is 2 atom % or more and is larger than the Mncontent of the first phase relative to the sum of Fe and Mn in the firstphase, or that both of these conditions are satisfied (i.e., the Mncontent of the second phase is twice or more the Mn content of the firstphase and is 2 atom % or more).

The Mn component content itself in the second phase does not exceed 100atom %. When the lower limit of the Mn content of the first phase is0.001 atom %, the Mn content of the second phase does not exceed 10⁵times the Mn content of the first phase. The Mn content of the secondphase is preferably 80 atom % or less of the Mn content of the firstphase. This is because when the Mn content exceeds 80 atom % (i.e., whenthe Mn content of the second phase exceeds 8×10⁴ times the Mn content ofthe first phase) while the second phase maintains the same crystalstructure as that of the first phase at ordinary temperature, thethermal stability of the whole magnetic material of the presentinvention may deteriorate.

In the above, the case described as the “Mn content” of the second phasebeing “twice or more” that of the first phase refers to a case in which,when the Mn content of each phase is calculated to one significantdigit, the Mn content of the second phase is twice or more the Mncontent of the first phase.

It is an objective of the present invention to lower coercive force byutilizing the magnetic anisotropy fluctuations in the above-mentionedrandom magnetic anisotropy model or a model in accordance therewith.Therefore, it is important either that the first phase and the secondphase, which are crystallographically independent, are magneticallycoupled at the nano level by exchange coupling, or that the Mn contentin the bcc phase including the first phase and the second phase has aspatial change at the nanoscale (this is sometimes referred to in thepresent invention as a “concentration fluctuation”). However, if the Mncomposition ratio of these two phases is too close, there are caseswhere the crystal orientations of the crystal phases are aligned in thesame direction, and in addition, the magnitude of the magnetocrystallineanisotropy constant is often smaller in the second phase. However, whenthe Mn content is less than 2 atom % relative to the sum of Fe and Mn inthe second phase, the crystalline magnetic anisotropy increases and whenaveraged the crystalline magnetic anisotropy value is not sufficientlysmall. As a result, a sufficiently low coercive force is not realized.Therefore, the preferable Mn content of the second phase is 2 atom % ormore relative to the sum of Fe and Mn in the second phase, and morepreferably 5 atom % or more. In the latter case, the magnetocrystallineanisotropy of the two phases is reduced to less than half that when Mnis not contained. When the Mn content is even more preferably 8 atom %or more, because this means that the magnetocrystalline anisotropy isvery small.

If there is a phase (first phase) in which the Mn content is lower thanthe Mn content of the whole magnetic material of the present invention,this means that in the same magnetic material there will be a phase(second phase) in which the Mn content is higher than that of themagnetic material of the present invention. Therefore, if isotropy isrealized as a result of those phases ferromagnetically coupling, thematerial will be the magnetic material of the present invention,specifically a soft magnetic material. Further, if the material isinterposed at the interface of the first phase, has a coercive forcewithin the appropriate range, and has an action of increasing electricresistance, then the material will be the magnetic material of thepresent invention, specifically a semi-hard magnetic material. Even whenthe material is not sufficiently isotropic, if there is spatialconcentration fluctuation of the Mn content in a given crystal phase,there will be fluctuation in the magnetic anisotropy, and the coerciveforce may decrease by a mechanism that is slightly different from therandom anisotropy model. In general, the magnetic material of thepresent invention, in which the coercive force decreases by such amechanism has a Mn content relative to the sum of Mn and Fe in themagnetic material of 10 atom % or less. The above is one characteristicof the magnetic material of the present invention that is not seen inmost existing soft magnetic materials such as electromagnetic steelsheets and sendust, which have highly homogenous compositions designedto thoroughly eliminate heterogeneity so as not to inhibit domain wallmovement. This characteristic can be said to be common with magneticmaterials in which magnetization reversal occurs due to the rotation ofmagnetization.

It may be noted that a state in which only the first phase or only thesecond phase is magnetically coupled at the nano level by exchangecoupling may be included in the present invention. Even in this case, itis important for the crystal axis directions of adjacent nanocrystalsnot to be aligned, and for the material to be isotropic or have ananoscale spatial distribution of the Mn content in the bcc phasecontaining the first phase and the second phase. However, in the presentinvention, it is impossible to achieve a magnetic material composed ofmicrocrystals of only the first phase or a magnetic material composed ofmicrocrystals of only the second phase, and even when such a structureis included, in the present invention, the first phase and the secondphase always exist in the magnetic material. The reason for this is thatthe formation of the nanocrystals per se plays a large role in thedisproportionation reaction in each of the processes of the reductionstep that kicks off with reduction of the nanoscale ferrite powdercontaining manganese that is used for producing the magnetic material ofthe present invention (in the present application, also referred to as“manganese ferrite nanopowder” or “Mn-ferrite nanopowder”). In thepresent application, a nanoscale ferrite powder is also referred to as a“ferrite nanopowder”, and the term “nanoscale” means, unless definedotherwise, a scale of 1 nm or more and less than 1 μm.

<Specification of Second Phase>

How to specify the second phase will now be described. First, asdescribed above, the first phase is an α-(Fe,Mn) phase, which is mainlyto guarantee a high saturation magnetization. The second phase is aphase whose Mn content relative to the sum of Fe and Mn contained inthat phase is higher than the Mn content relative to the sum of Fe andMn contained in the first phase. In the present invention, the secondphase may be an α-(Fe,Mn) phase whose Mn content is higher than the Mncontent of the whole magnetic material, or may be another crystal phase,an amorphous phase, or a mixed phase thereof. In any case, the softmagnetic material of the present invention has an effect of keeping thecoercive force low, so that even if a semi-hard magnetic material isincluded, there is an effect of imparting oxidation resistance andimproving electric resistivity. Therefore, since the second phase is anaggregate of phases having these effects, if the Mn content is higherthan that of the first phase, and it is possible to show the presence ofany of the phases exemplified above, that material can be understood asbeing the magnetic material of the present invention. If such a secondphase is not present and the material is composed only of the firstphase, that magnetic material will have poorer magnetic properties suchas coercive force, oxidation resistance, or electric conductivity, andinevitably processability will be poor and the molding steps will becomplicated.

If the second phase is an α-(Fe,Mn) phase, the Mn composition maycontinuously change from that of the first phase. Alternatively,depending on the method for identifying the material, it may appear asif the Mn composition of the second phase continuously changes from thefirst phase. In such a case as well, it is desirable that there is thefollowing compositional difference: the Mn content of the second phaseis larger than the Mn content of the first phase; and the Mn content ofthe second phase is further twice or more the Mn content of the firstphase and/or 2 atom % or more.

The composition ratio of Fe and Mn is, when the first phase and thesecond phase are combined, desirably 1:1 or less. In other words, the Mncontent relative to the total amount of Fe and Mn is desirably 0.01 atom% or more and 50 atom % or less.

The Mn content including the first phase and the second phase togetheris preferably 50 atom % or less in order to avoid a reduction in thesaturation magnetization, and is preferably 0.01 atom % or more in orderto avoid having no effect of adding the Mn on oxidation resistance andto avoid the coercive force becoming so high that it does not correspondto the intended use. Further, from the perspective of a good balancebetween oxidation resistance and magnetic properties, the Mn contentwhen the first phase and the second phase are combined is preferably0.02 atom % or more and 33 atom % or less, and a particularly preferablerange is 0.05 atom % or more and 25 atom % or less.

Although the volume ratio of the first phase and the second phase isarbitrary, the sum of the volume of the α-(Fe,Mn) phase in the firstphase, or the α-(Fe,Mn) phase in the first phase and the second phasebased on the whole magnetic material of the present invention includingthe first phase, the second phase, and the minor phase is preferably 5%by volume or more. Since the α-(Fe,Mn) phase is responsible for the mainmagnetization of the magnetic material of the present invention, theα-(Fe,Mn) phase volume is preferably 5% by volume or more in order toavoid a reduction in magnetization. Further, the α-(Fe,Mn) phase volumeis preferably 25% by volume or more, and more preferably 50% by volumeor more. To realize a particularly high magnetization without loweringthe electric resistivity too much, it is desirable to set the totalvolume of the α-(Fe,Mn) phases to 75% by volume or more.

In the second phase of the soft magnetic material of the presentinvention, it is preferable that there is a ferromagnetic phase or anantiferromagnetic phase (in the present application, weak magnetism isalso included therein), because there is an effect of reducing themagnetocrystalline anisotropy of the first phase.

<Example of Preferable Second Phase>

In the magnetic material of the present invention, as a representativeexample of the preferable second phase for ferromagnetism, firstly, a Mncontent of the second phase relative to a sum of Fe and Mn in the secondphase is higher than that of the first phase relative to a sum of Fe andMn in the first phase. Preferably, there is such an α-(Fe,Mn) phase thathas a Mn content of, relative to the sum of Fe and Mn in the secondphase, 0.1 atom % or more and 20 atom % or less, more preferably 2 atom% or more and 15 atom % or less, and particularly preferably 5 atom % ormore and 10 atom % or less.

A low coercive force is realized when the Mn content of the first phaserelative to the sum of Fe and Mn in the first phase is 5 atom % or moreand 10 atom % or less. However, when the Mn content is increased to sucha degree, a saturation magnetization of close to 2 T cannot beexhibited. Therefore, it is preferable to realize a magnetic materialhaving a large saturation magnetization and a small coercive force bycombining a first phase having a Mn content of less than 5 atom % and asecond phase having a Mn content of 5 atom % or more.

Next, examples of a preferable second phase may include an oxide phaseof both the Mn-ferrite phase and the wustite phase. The former isferromagnetic and the latter is antiferromagnetic, but either of themcan promote ferromagnetic coupling if it is in the first phase. Theseoxide phases may be nano-sized and very fine structures. In particular,for the wustite phase, it is several atomic layers thick and may befinely dispersed in the bcc phase or layered between the bccmicrocrystalline phases. When such an oxide layer is present, thecrystal orientation of the bcc phase may be uniform in the region ofseveral hundred nm to several tens of μm, but even with such amicrostructure, as long as the crystal grain size and the like is withinthe range of the present invention, such a material is considered to bethe magnetic material of the present invention. In particular, when amagnetic material having the above structure is a soft magneticmaterial, the coercive force is lowered by a mechanism slightlydifferent from the random anisotropy. It is assumed that the mechanismis as follows.

If, due to disproportionation, there is a difference between the Mncontent of the first phase relative to the sum of Fe and Mn in the firstphase and the Mn content of the second phase relative to the sum of Feand Mn in the second phase, and there is a concentration fluctuation inthe Mn content spatially at a fine nanoscale, spatial fluctuation of themagnetic anisotropic properties occurs. As a result, magnetizationreverses all of a sudden (as if a resonance phenomenon has occurred)when an external magnetic field is applied. The above concentrationfluctuation has the same effect on reducing coercive force not only whenthe second phase is an oxide phase, but also when it is an α-(Fe,Mn)phase.

Although examples in which the ferrite phase promotes ferromagneticcoupling are also known (in regards to this, see InternationalPublication No. WO 2009/057742 (hereinafter, referred to as “PatentDocument 1”), and N. Imaoka, Y. Koyama, T. Nakao, S. Nakaoka, T.Yamaguchi, E. Kakimoto, M. Tada, T. Nakagawa, and M. Abe, J. Appl.Phys., Vol. 103, No. 7 (2008) 07E129 (hereinafter, referred to as “NonPatent Document 3”)), in all of those cases, a ferrite phase is presentbetween Sm₂Fe₁₇N₃ phases of a hard magnetic material, and those phasesare ferromagnetically coupled to constitute an exchange spring magnet.

However, the present invention relates to a soft magnetic material and asemi-hard magnetic material, and exhibits completely different functionsfrom those of the above-mentioned hard magnetic exchange spring magnet.The present invention is the same in terms of the point of going throughan exchange interaction between first phases due to the presence of thesecond phase, which is a Mn-ferrite phase or a wustite phase. If such asecond phase is present so as to surround the first phase, electricresistance is also high, and coercive force is also reduced. Therefore,this is a particularly preferable second phase for the soft magneticmaterial of the present invention.

These two kinds of oxide phases are preferably 95% by volume or less ofthe whole magnetic material. This is because, for example, althoughMn-ferrite is a ferromagnetic material, its magnetization is lower thanthat of an α-(Fe,Mn) phase, and although wustite is also weakly magneticeven though it is antiferromagnetic, and hence there is somemagnetization, that magnetization is less than that of Mn-ferrite, sothat if the volume of either of these exceeds 95% by volume, themagnetization of the whole magnetic material may decrease. Morepreferably, the content of the oxide phase is 75% by volume or less, andparticularly preferably 50% by volume or less. In the case of producinga magnetic material having particularly high magnetization whilemaintaining electric resistivity to a certain extent, it is preferableto set the volume of the oxide phases to 25% by volume or less. On theother hand, when an oxide phase such as a wustite phase is present, theelectric resistivity increases. Therefore, when a wustite phase or thelike is intentionally contained for this reason, the volume fraction ispreferably 0.001% by volume or more. In order to have a wustite phaseand the like be present without excessively decreasing themagnetization, and to effectively improve the electric resistivity, thevolume is more preferably set to 0.01% by volume or more, andparticularly preferably to 0.1% by volume or more. Here, even when theoxide phases do not include Mn-ferrite and is assumed to be wustite, theabove-mentioned volume fraction range is the same.

As described above, an α-(Fe,Mn) phase having a larger Mn content thanthat of the first phase, a Mn-ferrite phase, and a wustite phase havebeen described as preferable examples of the second phase. These threephases may be a ferromagnetic material or an antiferromagnetic material.Therefore, if these phases are separated without ferromagnetic coupling,since the magnetic curve has additivity, the magnetic curves of thesemixed materials are simply the sum of the respective magnetic curves,and a smooth step is produced on the magnetic curve of the wholemagnetic material. For example, by observing the shape of the ¼ majorloop (the magnetic curve when swept from 7.2 MA/m to the zero magneticfield is called the ¼ major loop) of the magnetic curve of the wholemagnetic material, which is obtained by measuring the magnetization overa wide magnetic field range from 0 to 7.2 MA/m of an external magneticfield, it can be inferred that the smooth step on the ¼ major loop isdue to the above-mentioned circumstances or that there is certainly aninflection point based thereon. On the other hand, when these dissimilarmagnetic materials are formed as one body by ferromagnetic coupling, asmooth step or an inflection point is not seen on the major loop in therange of 7.2 MA/m to the zero magnetic field, but a monotonicallyincreasing magnetic curve with a convex portion at the top is produced.In order to estimate the existence of ferromagnetic coupling, inaddition to observing the fine structure at the grain boundary region asdescribed above, the above-mentioned detailed observation of themagnetic curve is also a measure.

Among the preferable second phases, which are the above-mentioned oxidephases, the wustite phase can be present stably even at a high reductiontemperature and molding temperature, and hence a wustite phase is a verypreferable phase in terms of forming the magnetic material of thepresent invention. Further, mainly in the reduction step, the α-(Fe,Mn)phases having various compositions caused by the disproportionationreaction from this phase are important phases responsible for supportingthe magnetic body manifested by the magnetic material of the presentinvention as the first phase or the first phase and the second phase. Inthe region where the Mn content is 0.5 atom % or more, because thereduction reaction progresses particularly via the wustite phase to ametal phase having high magnetism, in many cases, the α-(Fe,Mn) phasesare already directly ferromagnetically coupled with the wustite phasefrom the stage when those phases are produced by the disproportionationreaction. Therefore, those α-(Fe,Mn) phases are a very favorable phaseto utilize as the second phase of the magnetic material of the presentinvention, and in particular the soft magnetic material of the presentinvention.

<Composition Distribution>

In the examples of the present application, local composition analysisof the metal elements of the magnetic material of the present inventionis mainly carried out by EDX (energy dispersive X-ray spectroscopy), andthe composition analysis of the whole magnetic material is carried outby XRF (X-ray fluorescence elemental analysis). Generally, the Mncontent of the first phase and the second phase is measured by an EDXapparatus attached to an SEM (scanning electron microscope), an FE-SEM,a TEM (transmission electron microscope), or the like (in the presentapplication, this FE-SEM etc. equipped with an EDX is also referred toas an FE-SEM/EDX, for example). Depending on the resolution of theapparatus, if the crystal structure of the first phase and the secondphase is a fine structure of 300 nm or less, accurate compositionanalysis cannot be performed with an SEM or FE-SEM. However, to detectonly the difference in the Mn or Fe components of the magnetic materialof the present invention, those apparatus can be utilized in asupplementary manner. For example, in order to find a second phase thatis less than 300 nm and has a Mn content of 5 atom % or more, a certainpoint in the magnetic material is observed, and if the quantitativevalue of that point can be confirmed as having a Mn content of 5 atom %or more, then that means that a structure having a Mn content of 5 atom% or more or a part of such a structure is present within a diameter of300 nm centered on that one point. Conversely, to find a first phasehaving a Mn content of 2 atom % or less, a certain point is observed inthe magnetic material, and if the quantitative value of that point canbe confirmed as having a Mn content of 2 atom % or less, then that meansthat a structure having a Mn content of 2 atom % or less or a part ofsuch a structure is present within a diameter of 300 nm centered on thatone point.

When analyzing the composition by using the EDX apparatus provided withthe TEM, it is also possible to narrow down the electron beam to 0.2 nm,for example, and it is possible to perform very fine compositionanalysis. On the other hand, in order to investigate a certain areathoroughly and to gain an overall picture of the materials of thepresent invention, it is necessary to handle a large amount of data suchas, for example, 60 thousand points. In other words, it is necessary toappropriately select the composition distribution measurement methoddescribed above, and specify the compositional and structuralcharacteristics of the magnetic material of the present invention, suchas the composition of the first phase and the second phase, and thecrystal grain size.

Further, the composition of the α-(Fe,Mn) phase can also be determinedby confirming the diffraction peak positions by using an XRD (X-raydiffractometer). When the Mn content is 0 atom % or more and 3 atom % orless and 4 atom % or more, the diffraction peaks of the α-(Fe,Mn) phasetend to shift to a lower angle as the Mn content increases. Among thosepeaks, by observing the behavior of the (110) and (200) peaks, andcomparing with the diffraction positions of α-Fe (Fe-ferrite nanopowdernot including Mn or an M component was separately prepared and comparedas a comparative example while using the method of the present inventionas a reference) the Mn content of the α-(Fe,Mn) phase can be known toone significant digit. It is noted that in the region where the Mncontent is more than about 3 atom % and less than 4 atom %, thediffraction angle tends to conversely increase with increasing Mncontent, whereas at 4 atom %, the diffraction angle is higher than thatwhen the Mn content is 0%, namely, the lattice constant is smaller thanin the α-Fe phase.

Further, by comprehensively combining these composition analysismethods, namely, XRD, FE-SEM, TEM, or the like, it is possible to knowthe orientation and composition distribution of the crystal grains. TheMn composition, which is a characteristic of the present invention, isuseful for verifying whether various crystal phases are present due todisproportionation, and whether their crystal axes are randomly orientedor not. Furthermore, to distinguish the α-(Fe,Mn) phase from the otheroxide phases, such as a wustite phase, it is convenient and effective toanalyze the oxygen characteristic X-ray surface distribution map using,for example, SEM-EDX.

<Composition of Whole Magnetic Material>

The composition of the whole magnetic material in the present inventionis in the range of, based on the composition of the whole magneticmaterial, 20 atom % or more and 99.999 atom % or less of the Fecomponent, 0.001 atom % or more and 50 atom % or less of the Mncomponent, and 0 atom % or more and 55 atom % or less of O (oxygen).Preferably, all of these ranges are simultaneously satisfied. Further,an alkali metal may be contained in the range of 0.0001 atom % or moreand 5 atom % or less. It is desirable that the minor phase including Kand the like does not exceed 50% by volume of the whole.

It is preferable that Fe is 20 atom % or more, because a reduction inthe saturation magnetization can be avoided. It is preferable that Fe is99.999 atom % or less, because a reduction in the oxidation resistanceand deterioration in workability can be avoided. Also, it is preferablethat the Mn component is 0.001 atom % or more, because a reduction inthe oxidation resistance and deterioration in workability can beavoided. It is preferable that the Mn component is 50 atom % or less,because a reduction in the saturation magnetization can be avoided. WhenO is an important element for forming the second phase, it is preferablethat O is in a range of 55 atom % or less, because not only a reductionin the saturation magnetization can be avoided, but a situation in whichthe disproportionation reaction in the first phase and the second phaseby reduction of the manganese ferrite nanopowder does not occur, makingit more difficult to develop to a low coercive force soft magneticmaterial can be avoided. Although the magnetic material of the presentinvention does not necessarily need to contain oxygen, it is desirablethat even a slight amount be contained in order to obtain a magneticmaterial with remarkably high oxidation resistance and electricresistivity. For example, it is possible to passivate the surface of themetal powder reduced by the gradual oxidation step (described later), orto cause an oxide layer composed of wustite and the like to be presentat a part of the crystal grain boundary of the solid magnetic materialby that passivation action. In this case, the respective compositionranges of the whole magnetic material of the present invention aredesirably 20 atom % or more and 99.998 atom % or less of the Fecomponent, 0.001 atom % or more and 50 atom % or less of the Mncomponent, and 0.001 atom % or more and 55 atom % or less of O.

A more preferable composition of the magnetic material of the presentinvention is 50 atom % or more and 99.98 atom % or less of the Fecomponent, 0.01 atom % or more and 49.99 atom % or less of the Mncomponent, and 0.01 atom % or more and 49.99 atom % or less of O. Inthis range, the magnetic material of the present invention has a goodbalance between saturation magnetization and oxidation resistance.

Furthermore, the magnetic material of the present invention having acomposition in which the Fe component is in the range of 66.95 atom % ormore and 99.9 atom % or less, the Mn component is in the range of 0.05atom % or more and 33 atom % or less, and O is in the range of 0.05 atom% or more and 33 atom % or less is preferable from the perspective ofhaving excellent electromagnetic properties and excellent oxidationresistance.

Within the above composition ranges, when the magnetic material of thepresent invention is to have an excellent performance, in particular, amagnetization of 1 T or more, a preferable composition range is 79.95atom % or more and 99.9 atom % or less for the Fe component, 0.05 atom %or more and 20 atom % or less for the Mn component, and 0.05 atom % ormore and 20 atom % or less of O.

Since it also depends on the Mn component content, and hence cannot beunconditionally stated, in the present invention the semi-hard magneticmaterial tends to contain more oxygen than the soft magnetic material.

<Magnetic Properties, Electrical Properties, and Oxidation Resistance>

One aspect of the present invention is a magnetic material havingmagnetic properties, electrical properties, and oxidation resistancesuitable for soft magnetic applications with a coercive force of 800 A/mor less. These points are now described below.

The term “magnetic properties” as used herein refers to at least one ofthe magnetic material's magnetization J (T), saturation magnetizationJ_(s) (T), magnetic flux density (B), residual magnetic flux densityB_(r) (T), exchange stiffness constant A (J/m), magnetocrystallineanisotropy magnetic field Ha (A/m), magnetocrystalline anisotropy energyEa (J/m³), magnetocrystalline anisotropy constant K₁ (J/m³), coerciveforce H_(cB) (A/m), intrinsic coercive force H_(cJ) (A/m), permeabilityμμ₀, relative permeability μ, complex permeability μ_(r)μ₀, complexrelative permeability μ_(r), its real term μ′, imaginary term μ″, andabsolute value |μ_(r)|. In the present specification, A/m from the SIunit system and Oe from the cgs Gauss unit system are both used as theunits of the “magnetic field”. The formula for conversing between thosevalues is 1 (Oe)=1/(4π)×10³ (A/m). More specifically, 1 Oe is equivalentto about 80 A/m. As the units for the “saturation magnetization” and“residual magnetic flux density” in the present specification, T fromthe SI unit system and emu/g from the cgs Gauss unit system are bothused. The formula for converting between those values is 1(emu/g)=4π×d/10⁴ (T), where d (Mg/m³=g/cm³) represents density.Therefore, since d=7.87 for Fe, Fe having a saturation magnetization of218 emu/g has a saturation magnetization value M_(s) in the SI unitsystem of 2.16 T. In the present specification, unless stated otherwise,the term “coercive force” refers to the intrinsic coercive force Ha.

The term “electrical properties” used herein refers to the electricresistivity (=volume resistivity) ρ (Ωm) of the material. The term“oxidation resistance” used herein refers to a change over time in themagnetic properties in various oxidizing atmospheres, for example, aroom-temperature air atmosphere.

The above-mentioned magnetic properties and electrical properties arecollectively referred to as “electromagnetic properties”.

In the magnetic material of the present invention, it is preferable thatthe magnetization, the saturation magnetization, the magnetic fluxdensity, the residual magnetic flux density, and the electricresistivity are higher. For the saturation magnetization, a value ashigh as 0.3 T or 30 emu/g or more is desirable. For soft magneticmaterials in particular, a value as high as 100 emu/g or more isdesirable. For electric resistivity, a value as high as 1.5 μΩm or moreis desirable. Other magnetic properties of the present invention, suchas the magnetocrystalline anisotropy constant, the coercive force, thepermeability, the relative permeability, and the like are appropriatelycontrolled depending on the application and based on the material is tobe formed as a semi-hard magnetic material or as a soft magneticmaterial. In particular, depending on the application, the permeabilityand relative permeability do not always have to be high. As long as thecoercive force is sufficiently low and the iron loss is suppressed to alow level, for example, the relative permeability may even be adjustedto a magnitude in the range of 10⁰ to around 10⁴. In particular, bysuppressing the magnetic saturation under a direct-current superimposedmagnetic field, it is possible to suppress the deterioration inefficiency and facilitate linear control, or based on the relationalexpression (1), each time the permeability is reduced by one digit, thecritical thickness at which eddy current loss occurs can be increased bya factor of about 3.2. One of the characteristics of the presentinvention lies in comprising a magnetization reversal mechanism that isbased mainly on direct rotation of magnetization, and not only onmagnetization reversal due to domain wall movement. As a result, thecoercive force is low, eddy current loss due to domain wall movement issmall, and iron loss can be suppressed to a low level. Moreover, it ispossible to generate some local magnetic anisotropy at the crystalboundary for suppressing magnetization rotation by the external magneticfield, and to reduce permeability.

Incidentally, in the present invention, the reason why such permeabilityadjustment is possible is that because the electric resistivity of themagnetic material is large even when sintered as is, iron loss due tothe eddy current is small, and therefore even if the hysteresis lossincreases a little due to designing the material to suppresspermeability by sacrificing a little bit of coercive force, the totaliron loss can be kept at a low level.

The soft magnetic material of the present invention exhibits an electricresistivity of 1.5 μΩm or more, and in semi-hard magnetic materials, itexhibits even higher electric resistivity.

In the soft magnetic material of the present invention exhibiting anelectric resistivity of 10 μΩm or more, since the saturationmagnetization tends to decrease as the electric resistivity increases,it is necessary to determine the composition and degree of reduction ofthe raw materials according to the desired electromagnetic properties.In particular, an electric resistivity of less than 1000 μΩm ispreferable for obtaining the characteristic that the magnetization ofthe magnetic material of the present invention is high. Therefore, thepreferred range of electric resistivity is 1.5 μΩm or more and 1000 μΩmor less.

<Crystal Boundaries>

As described above, whether the magnetic material of the presentinvention becomes soft magnetic or semi-hard magnetic depends on themagnitude of the coercive force, but in particular it is closely relatedto the fine structure of the magnetic material. Although an α-(Fe,Mn)phase may at a glance look as if they are a continuous phase, as shownin FIG. 1, the magnetic material contains many heterogenous phaseinterfaces and crystal grain boundaries. Further, the magnetic materialcontains crystals such as twin crystals including simple twins such ascontact twins and penetrating twins, recurring twins such aspolysynthetic twins, cyclic twins, and multiple twins, intergrowths, andskeleton crystals (in the present invention, when crystals areclassified not only by the heterogenous phase interface and thepolycrystalline grain boundary but also by the various crystal habits,tracht, intergrowth structures, dislocations, and the like describedabove, those boundary surfaces are collectively referred to as “crystalboundaries”). In many cases, unlike linear grain boundaries which aregenerally often seen, the crystal boundaries are often presented as agroup of curves, and furthermore, in such a structure, there is a largedifference in Mn content depending on location. The magnetic material ofthe present invention having such a fine structure is often a softmagnetic material.

In the case where the magnetic material of the present invention is asoft magnetic material, when the second phase is an α-(Fe,Mn) phase,starting from a manganese ferrite nanopowder, as the first phase and thesecond phase undergo grain growth, and as the reduction reactionprogresses, the oxygen in the crystal lattice is lost in conjunctionwith the disproportionation reaction of the composition, eventuallycausing a large reduction in volume of normally up to 52% by volume. Asa result of this, the first phase and the second phase, which areα-(Fe,Mn) phases, have diverse microstructures, such as crystals thatare observed in precious stones such as quartz and minerals and rockssuch as pyrite and aragonite, that are in a reduced form on a nanoscaleand contain various phases and nanocrystals with various Mn contents intheir interior.

The structures seen at the grain boundaries and in intergrowths may alsoexhibit a difference in Mn content depending on the observed location,and hence are a heterogenous phase interface in some cases.

When Mn is added to Fe, the magnetocrystalline anisotropy energydecreases. Further, even if up to about 8 atom % of Mn is added to Fe,the magnetization does not decrease so much, and is only about a 5%reduction. In addition, if the amount of Mn is more than 0 and less than1 atom %, depending on the conditions, the saturation magnetization perunit mass may exceed iron. The reason why the saturation magnetizationper unit mass improves only in this region is unknown, but it is thoughtto be due to the crystal phases of the first phase of the nanopowder orthe crystal phases of the first phase and the second phase of thenanopowder being stacked at the nanoscale like, for example, (shapemagnetism anisotropy averaged) pearlite, and/or due to microcrystalsbeing bonded at the nanoscale like a eutectic structure, causing theferromagnetic phase lattice to become distorted or stretched, and it ispresumed that this is an effect not to be seen with macroscopiccrystals. Generally, it is known that when Mn is added to Fe, which is amacroscopic bulk crystal, in a dilute region, the influence of the Mn onthe saturation magnetic moment follows almost simple dilution, and likeelements such as Co, Ni, Ir, Pt, and Rh, it is thought that there is noeffect of increasing the saturation magnetic moment of the Fe alloy inthe dilute region from this addition.

Further, addition of 2 atom % or more of Mn has the effect ofmaintaining the microstructure having composition fluctuation even whenthe cooling rate is slow, and even in the region of more than 0 and lessthan 2 atom %, a fine structure tends to be maintained if the coolingcan be controlled to within a certain cooling rate.

<Random Magnetic Anisotropy Model and Coercive Force Reduction MechanismUnique to Present Invention>

For the soft magnetic material of the present invention described by therandom anisotropy model, it is important that the following threeconditions are satisfied.

(1) Crystal grain size of the α-(Fe,Mn) phase is small.(2) There is ferromagnetic coupling by exchange interaction.(3) There is random orientation.

Item (3) is not always necessary, when the Mn content in the bcc phaseis 10 atom % or less. In this case, the reduction in coercive forceoccurs based on a different principle from the random anisotropy model.Specifically, magnetic anisotropy fluctuations occur based onconcentration fluctuations in the nanoscale Mn content due tointeractions between any one or more of the first phase and the secondphase, the first phases themselves, or the second phases themselves.This promotes magnetization reversal, and the coercive force is reduced.The magnetization reversal mechanism based on this mechanism is uniqueto the present invention, and was discovered for the first time by thepresent inventors as far as the inventors are aware.

In cases where the grain growth during reduction is insufficient, orwhere the grains do not fuse with each other so as to form a continuousferromagnetic phase, or where there is phase separation in which grainsseparate without going through a phase which is not ferromagnetic orantiferromagnetic, to bring the coercive force of the magnetic materialpowder of the present invention into the soft magnetic region, it isdesirable to subsequently solidify the magnetic material by sintering orthe like, namely, form “state in which the first phase and the secondphase are continuously bonded to each other directly or via a metalphase or an inorganic phase to form a massive state as a whole”.

In order to achieve the above item (2), namely, ferromagnetic couplingby exchange interaction, since the exchange interaction is aninteraction or force that acts within a short range in the order ofseveral nm, when first phases are coupled to each other, the phases aredirectly bonded, and when a first phase and a second phase or secondphases are coupled to each other, it is necessary for the second phaseto be ferromagnetic or antiferromagnetic in order to transmit theexchange interaction. Even if a part of the first phase and/or thesecond phase is in a superparamagnetic region, since the material itselfis ferromagnetic or antiferromagnetic in the bulk state, as long as thesurrounding ferromagnetic or antiferromagnetic phase is sufficientlyexchange coupled, those phases may be able to transmit an exchangeinteraction.

In the case of the semi-hard magnetic material of the present invention,although not limited to the above, in order to obtain a semi-hardmagnetic material having a high residual magnetic flux density, theabove-mentioned solidification is necessary.

<Average Crystal Grain Size of First Phase, Second Phase, and WholeMagnetic Material>

The average crystal grain size of the first phase or the second phase ofthe soft magnetic material of the present invention or the averagecrystal grain size of the whole magnetic material is preferably 10 μm orless. When the average crystal grain size of the first phase and thesecond phase is 10 μm or less, the average crystal grain size of thewhole magnetic material is 10 μm or less.

In particular, regarding the soft magnetic material of the presentinvention, in order to realize a low coercive force by the above randommagnetic anisotropy model or the mechanism unique to the presentinvention, it is preferable that either the first phase or the secondphase is in the nano region. When both the first phase and the secondphase are ferromagnetic phases, it is preferable that the averagecrystal grain size of both phases is 10 μm or less, and preferably lessthan 1 μm in order to realize a low coercive force based on the randommagnetic anisotropy model. This average crystal grain size is morepreferably 500 nm or less, and particularly preferably 200 nm or less,because a remarkable reduction effect of the coercive force by themechanism unique to the present invention can be realized, although thisdoes depend on the Mn content as well. In the above case, since the K₁of the first phase is larger than the second phase in many cases,particularly when the first phase is 10 μm or less, preferably 500 nm orless, and more preferably 200 nm or less, the coercive force becomesvery small, and a soft magnetic material suitable for varioustransformers, motors, and the like is obtained.

On the other hand, if this average crystal grain size is less than 1 nm,superparamagnetism occurs at room temperature, and magnetization andpermeability may become extremely small. Therefore, it is preferablethat this average crystal grain size is 1 nm or more. As mentionedabove, if crystal grains smaller than 1 nm or amorphous phases arepresent, these need to be sufficiently coupled to crystal grains of 1 nmor more in size by exchange interaction.

Further, when the second phase is not a ferromagnetic phase, the secondphase is not involved in reducing the coercive force by the randomanisotropy model or the mechanism unique to the present invention, butits presence increases the electric resistivity, and hence it ispreferable for that component to be present.

In the case of the semi-hard magnetic material of the present invention,in order to express a coercive force, it is effective to maintain theaverage crystal grain size of the first phase at the nano level, andeither employ a suitable surface oxide layer as the second phase orcause a second phase having an average crystal grain size of several nmto be present at the grain boundary of the first phase to therebymaintain a high magnetization and impart oxidation resistance whilemaintaining the coercive force of the semi-hard magnetic region.

<Measurement of Crystal Grain Size>

Measurement of the crystal grain size of the present invention iscarried out using an image obtained by SEM, TEM, or metallographicmicroscopy. The crystal grain size is obtained by, within an observedrange, observing not only the heterogenous phase interfaces and crystalgrain boundaries but all the crystal boundaries, and taking the diameterof the crystal region of the surrounded portion to be the crystal grainsize. When the crystal boundary is difficult to see, the crystalboundary may be etched by a wet method using a Nital solution or thelike, a dry etching method, or the like. The average crystal grain sizeis, in principle, obtained by selecting a representative portion andmeasuring a region containing at least 100 crystal grains. Although thenumber of grains may be less than this, in that case the measurementneeds to be carried out on a portion that is statistically sufficientlyrepresentative of the whole. The average crystal grain size is obtainedby photographing the observation area, defining an appropriaterectangular quadrilateral area on the photographic plane (the enlargedprojection plane on the target photographic plane), and applying theJeffry method to the interior of that defined area. When observing by anSEM or a metallurgical microscope, the crystal boundary width may be toosmall in relation to the resolution and may not be observed, but in thatcase the measured value of the average crystal grain size gives theupper limit of the actual crystal grain size. Specifically, it is enoughif the average crystal grain size measurement value has an upper limitof 10 μm. However, there is a possibility that part or all of themagnetic material may be below 1 nm, which is the lower limit of thecrystal grain size, due to phenomena such as having no clear diffractionpeaks in XRD and superparamagnetism being confirmed on the magneticcurve. In such a case, the actual crystal grain size must be determinedagain by TEM observation.

In addition, in the present invention, it is sometimes necessary tomeasure the crystal grain size not in relation to the crystal boundary.More specifically, for example, when the crystal structure is finelymodulated due to the concentration fluctuations of the Mn content, thecrystal grain size of the magnetic material of the present inventionhaving such a fine structure is the modulation width of that Mn content.The method for determining the crystal grain size is specificallyillustrated in the methods using TEM-EDX analysis in Examples 20 and 21of the present invention.

<Measurement of Crystallite Size>

In the present invention, phase separation occurs due to thedisproportionation reaction, and a composition width occurs in the Mncontent of the bcc phase of the first phase and/or the second phase.Since the X-ray diffraction line peak positions vary depending on the Mncontent, even if the line width of the (200) diffraction line of the bccphase is calculated, for example, this is fairly meaningless in terms ofdetermining the crystallite size. Here, the term “crystallite” refers toa small single crystal at the microscopic level forming the crystalsubstance, which is smaller than the individual crystals (so-calledcrystal grains) forming the polycrystal.

On the other hand, when the Mn content of the bcc phase is up to 1 atom%, the deviation of the (200) diffraction line is within 0.026 degrees(Co-Kα line). Therefore, in the range of 1 nm or more and less than 100nm, it is meaningful to measure the crystallite size to one significantdigit.

In the present invention, the crystallite size of the bcc phase wascalculated by using the (200) diffraction line width excluding theinfluence of the Km diffraction line and the Scherrer equation, andtaking the dimensionless shape factor to 0.9.

The bcc phase may be a phase in which at least the first phase has thebcc phase (i.e., a case in which only the first phase has the bcc phaseand a case in which both the first phase and the second phase have thebcc phase), but a preferable crystallite size range of the bcc phase is1 nm or more and less than 100 nm.

When the crystallite size is less than 1 nm, superparamagnetism occursat room temperature, and magnetization and permeability may becomeextremely small. Therefore, it is preferable that this crystallite sizeis 1 nm or more.

The crystallite size of the bcc phase is preferably less than 100 nm,because the coercive force enters the soft magnetic region and becomesextremely small, and a soft magnetic material suitable for varioustransformers, motors, and the like is obtained. Further, at 50 nm orless, not only a high magnetization exceeding 2 T, which is a low regionof the Mn content, can be obtained, but also a low coercive force can beachieved at the same time, and hence this is a very preferable range.

<Size of Soft Magnetic Material>

The size of the powder of the soft magnetic material of the presentinvention is preferably 10 nm or more and 5 mm or less. If this size isless than 10 nm, the coercive force does not become sufficiently small,and if the size exceeds 5 mm, a large strain is applied duringsintering, and the coercive force conversely increases unless there isan annealing treatment after solidification. More preferably the size is100 nm or more and 1 mm or less, and particularly preferably is 0.5 μmor more and 500 μm or less. If the average powder particle diameter iscontained in this region, a soft magnetic material with a low coerciveforce is obtained. In addition, the particle size distribution ispreferably sufficiently wide within each average powder particlediameter range defined above, because high filling is easily achievedwith a relatively small pressure and the magnetization based on thevolume of the solidified molded body is increased. When the powderparticle diameter is too large, movement of the domain walls may beexcited, and due to the heterogenous phases formed by thedisproportionation reaction in the production process of the softmagnetic material of the present invention, that domain wall movement ishindered, which can conversely result in the coercive force becominglarger. Therefore, when molding the soft magnetic material of thepresent invention, it can be better for the surface of the magneticmaterial powder of the present invention having an appropriate powderparticle diameter to be in an oxidized state. In the alloy containing Mnof the present invention, since the structure is finer due to thedisproportionation reduction reaction, even if the surface is oxidizedto some extent by oxidation, in many cases there is no large influenceon the magnetization rotation of the interior portion, and the oxidationresistance is extremely high. In addition, even if the free energy ofoxidation is considerably lower than that of Fe and the surface isoxidized by external oxygen, Mn is preferentially oxidized and hence aneffect of suppressing oxidation of the Fe component, which forms themain part of the magnetic material of the present invention, is assumedto be also caused. Therefore, depending on the composition, shape, andsize of the magnetic material powder of the present invention,performing appropriate gradual oxidation of the powder surface, carryingout each step in air, and performing the solidification treatment in aninert gas atmosphere or the like rather than in a reducing atmosphereare also effective for stabilizing the coercive force.

<Size of Semi-Hard Magnetic Material>

The magnetic powder of the semi-hard magnetic material of the presentinvention is preferably in a range of 10 nm or more and 10 μm or less.If this average powder particle diameter is less than 10 nm, molding ishard to carry out, and when the magnetic material is used dispersed in asynthetic resin or ceramic, dispersibility is very poor. In addition, ifthe powder particle diameter exceeds 10 μm, since the coercive forcereaches the soft magnetic region, the magnetic material falls into thecategory of a soft magnetic material of the present invention. A morepreferable powder particle diameter is 10 nm or more and 1 μm or less.Within this range, the magnetic material is a semi-hard magneticmaterial with a good balance between saturation magnetization andcoercive force.

<Measurement of Average Powder Particle Diameter>

The powder particle diameter of the magnetic material of the presentinvention is mainly evaluated based on its median diameter calculatedfrom a distribution curve obtained by measuring the volume-equivalentdiameter distribution using a laser diffraction type particle sizedistribution meter. Alternatively, the powder particle diameter may alsobe calculated by choosing a photograph of the powder obtained by SEM orTEM, or a representative portion based on a metallographic micrograph,measuring the diameter of at least 100 particles, and volume-averagingthe diameters of those particles. Although the number of particles maybe less than this, in that case the measurement needs to be carried outon a portion that is statistically sufficiently representative of thewhole. In particular, when measuring the particle size of a powdersmaller than 500 nm or a powder exceeding 1 mm, priority is given to amethod using SEM or TEM. In addition, when a total number ofmeasurements n is performed using N types (N≤2) of measurement method ormeasurement apparatus in combination (N≤n), the numerical values R_(n)thereof needs to be within a range of R/2≤R_(n)≤2R. In that case, thepowder particle diameter is determined based on R, which is thegeometric average of the lower limit and the upper limit.

As described above, in principle, the powder particle diameter of themagnetic material of the present invention is measured by (1) when themeasurement value is 500 nm or more and 1 mm or less, preferentiallyusing the laser diffraction type particle size distribution meter, (2)when the measurement value is less than 500 nm or more than 1 mm,preferentially using microscopy, and (3) when the measurement value is500 nm or more and 1 mm or less and methods (1) and (2) are to becombined, calculating the average powder particle diameter by using theabove-mentioned R. In the present application, the powder particlediameter is expressed to one to two significant digits in the case ofmethods (1) or (2), and in the case of (3) is expressed to onesignificant digit. The reason why the methods for measuring the powderparticle diameter are used together is that when the powder particlediameter is just above 500 nm or just below 1 mm, there is a possibilitythat with method (1) an inaccurate value is obtained even when expressedto one significant digit, while on the other hand, for method (2), ittakes time and effort to confirm that the measurement value is not localinformation. Therefore, it is very rational to first obtain the value ofthe average powder particle diameter by method (1), then obtain thevalue easily by method (2), comparatively look at the two values anddetermine the average powder particle diameter by using theabove-mentioned R. In the present application, the average particlediameter of the powder of the magnetic material of the present inventionis determined by the above method. However, if methods (1) and (3), ormethods (2) and (3) do not match to one significant digit, R must bedetermined by precisely measuring using method (1) or (2) again based onthe average powder particle diameter range. However, when there areobvious inappropriate reasons, such as when there is clearly strongagglomeration and it would be inappropriate to determine the powderparticle diameter by method (1), or when the powder is too uneven andthe powder particle diameter estimated from the sample image is clearlydifferent and it would be inappropriate to determine the powder particlediameter by method (2), or when due to the specification of themeasurement apparatus, classifying based on a size of 500 nm or 1 mm asthe standard for determining the powder particle diameter measurementwould be inappropriate, it is acceptable to disregard the aboveprinciple and re-select one of the methods (1), (2) or (3) for thatparticular case. Specifically, within the scope of the measurementmethods (1) to (3), the most appropriate method for obtaining the volumeaverage value of the powder particle diameter as close as possible tothe true value may be selected by grasping capturing the true form ofthe magnetic material. If it is only necessary to distinguish themagnetic material of the present invention from other magneticmaterials, it is sufficient for the average powder particle diameter tobe determined to one significant digit.

For example, in the case of reducing a manganese ferrite nanopowderhaving a Mn content of 10 atom % or less at 900° C. or more, themacroscopic powder shape is a three-dimensional network structure inwhich many hollow portions, which are through-holes, are containedinside, and hence the powder may become sponge-like. These hollowportions are thought to be formed by large volume reductions caused byoxygen leaving the crystal lattice and the Mn component subliming asgrain growth progresses in the reduction reaction. The powder particlediameter in this case is measured including the volume of the interiorhollow portions.

<Solid Magnetic Material>

The magnetic material of the present invention can be used as a magneticmaterial in which the first phase and the second phase are continuouslybonded to each other directly or via a metal phase or an inorganic phaseto form a massive structure as a whole (in the present application, alsoreferred to as “solid magnetic material”). Further, as described above,when many nanocrystals are already bonded in the powder, the powder maybe molded by mixing with an organic compound such as a resin, aninorganic compound such as glass or ceramic, a composite materialthereof or the like.

<Packing Factor>

The packing factor is not particularly limited as long as the objects ofthe present invention can be achieved. However, when the magneticmaterial of the present invention contains a small amount of the Mncomponent, from the perspective of a balance among oxidation resistance,electric resistivity, and magnetization level, it is preferable to setthe packing factor to 60% by volume or more and 100% by volume or less.

As used herein, the term “packing factor” refers to the ratio, expressedas a percentage, of the volume of the magnetic material of the presentinvention relative to the volume of the whole magnetic material of thepresent invention including voids (i.e., volume occupied only by themagnetic material of the present invention, excluding the portion thatis not the magnetic material of the present invention, such as voids andresin).

A more preferable range of the packing factor is 80% or more, andparticularly preferable is 90% or more. Although the magnetic materialof the present invention has high oxidation resistance to begin with, asthe packing factor is increased, the oxidation resistance furtherincreases, and there is a wider range of applications that the magneticmaterial of the present invention can be applied to. In addition, thesaturation magnetization is also improved, and a high performancemagnetic material can be obtained. Further, for the soft magneticmaterial of the present invention, there are also the effects of anincrease in bonding between the powder particles and a reduction in thecoercive force.

<Characteristics of Magnetic Powder and Solid Magnetic Material ofPresent Invention>

One of the major characteristics of the magnetic material powder of thepresent invention is that it is a sinterable powder material likeferrite. Various solid magnetic materials having a thickness of 0.5 mmor more can easily be produced. Even various solid magnetic materialshaving a thickness of 1 mm or more, and even 5 mm or more, can beproduced comparatively easily by sintering or the like as long as thethickness is 10 cm or less.

Further, one of the characteristics of the magnetic material of thepresent invention is that it has a large electric resistivity. Whileother metallic rolled materials and thin ribbon materials are made by amanufacturing method in which crystal grain boundaries, heterogenousphases, or defects are not included, the magnetic material powder of thepresent invention includes many crystal boundaries and various phases,and it itself has an effect of increasing the electric resistivity. Inaddition, when solidifying the powder, because in particular a surfaceoxidized layer of the powder before solidification (i.e., layer having ahigh oxygen content that is present on the surface of the first phase orthe second phase, such as a MnO, wustite, magnetite, Mn-ferrite,ilmenite, titanohematite, or amorphous layer, and among those, an oxidelayer containing a large amount of Mn) and/or a metal layer (i.e., ametal layer containing a large amount of Mn) is intercalated, theelectric resistivity of the bulk body also increases.

In particular, examples of preferable constituent compounds of a surfaceoxide layer for increasing the electric resistivity include at least oneof MnO, wustite, and Mn-ferrite.

The reason why the magnetic material of the present invention has theabove characteristics is that the present invention mainly provides abuild-up type bulk magnetic material by producing a magnetic materialthat has a high magnetization and that is formed by a method which isessentially different from other metallic soft magnetic materials forhigh frequency applications, namely, by first producing a metal powderhaving nanocrystals by reducing a manganese ferrite nanopowder and thenforming a solid magnetic material by molding the magnetic powder.

Further, as described above, since the electric resistance is higherthan existing metal-based soft magnetic materials represented by siliconsteel, for example, it is possible to considerably simplify thelamination step and the like that are normally required when producingrotating devices and the like. Assuming that the electric resistivity ofthe magnetic material of the present invention is about 30 times that ofsilicon steel, the limit of the thickness at which an eddy current doesnot occur is, based on the relational expression (1), about 5 times thethickness. Therefore, even when lamination is required, the number oflaminations is also ⅕. For example, even when applied to a stator of amotor in a high rotation range with a frequency of 667 Hz, a thicknessof 1.5 mm is permitted.

For example, when the solid magnetic material of the present inventionis to be applied as a soft magnetic material as described above, thesolid magnetic material may be used in a wide variety of shapes inaccordance with the application.

The solid magnetic material of the present invention does not contain abinder such as a resin, has high density, and can be easily processedinto an arbitrary shape by an ordinary processing machine by cuttingand/or plastic working. In particular, one of the major characteristicsof the solid magnetic material is that it can be easily processed into aprismatic shape, a cylindrical shape, a ring shape, a disk shape, a flatsheet shape or the like having high industrial utility value. It is alsopossible to process the solid magnetic material into those shapes andthen subject to cutting and the like for processing into a roof tileshape or a prismatic shape having an arbitrary base shape. Specifically,it is possible to easily perform cutting and/or plastic working into anarbitrary shape or any form surrounded by flat surfaces or curvedsurfaces, including cylindrical surfaces. Here, the term “cutting”refers to cutting general metal materials. Examples include machineprocessing by a saw, a lathe, a milling machine, a drilling machine, agrinding stone, and the like. The term “plastic working” refers to aprocess such as die cutting by a press, molding, rolling, explosionforming, and the like. Further, in order to remove distortion after coldworking, annealing can be performed in the range of ordinary temperatureto 1290° C.

<Production Method>

Next, the method for producing the magnetic material of the presentinvention will be described, but the present invention is notparticularly limited thereto.

The method for producing the magnetic material of the present inventioncomprises:

(1) a manganese ferrite nanopowder production step; and(2) a reduction step,

and may optionally further comprise any one or more of the followingsteps:

(3) a gradual oxidation step;(4) a molding step; and(5) an annealing step.

Each step is now described in more detail.

(1) Manganese Ferrite Nanopowder Production Step (in the presentapplication, also referred to as “step (1)”)

Examples of a preferable step of producing the nanomagnetic powder,which is a raw material of the magnetic material of the presentinvention, include a method of synthesizing at room temperature using awet synthesis method.

Examples of known methods for producing a ferrite fine powder include adry bead mill method, a dry jet mill method, a plasma jet method, an arcmethod, an ultrasonic spray method, an iron carbonyl vapor phasecracking, and the like. Any of these methods is a preferable productionmethod, as long as the magnetic material of the present invention isformed. However, to obtain nanocrystals having a disproportionatedcomposition, which is the essence of the present invention, it ispreferable to mainly employ a wet method using an aqueous solutionbecause it is the simplest.

This production step is carried out by applying the “ferrite platingmethod” described in Patent Document 1 to the step for producing themanganese ferrite nanopowder used for producing the magnetic material ofthe present invention.

The ordinary “ferrite plating method” is applied not only to powdersurface plating but also to thin films and the like. The reactionmechanism and the like of the ferrite plating method have already beendisclosed (e.g., see Masaki Abe, Journal of the Magnetics Society ofJapan, Volume 22, No. 9 (1998), page 1225 (hereinafter, referred to as“Non Patent Document 4”) and International Publication No. WO2003/015109 (hereinafter, referred to as “Patent Document 2”)). However,unlike such a “ferrite plating method”, in this production step, thepowder surface, which serves as the base material of the plating, is notused. In this production step, the raw materials (e.g., manganesechloride and iron chloride) used for ferrite plating are reacted insolution at 100° C. or less to directly synthesize the ferrous andcrystalline manganese ferrite nanopowder itself. In the presentapplication, this step (or method) is referred to as “manganese ferritenanopowder production step” (or “manganese ferrite nanopowder productionmethod”).

A “manganese ferrite nanopowder production step” in which the nanopowderhas a spinel structure is described below as an example.

An appropriate amount of an aqueous solution adjusted in advance to anacidic region is placed in a container (in the present application, alsoreferred to as a reaction field), and while subjecting to ultrasonicwave excitation at room temperature under atmospheric pressure ormechanical stirring at an appropriate strength or revolution number, apH adjusting solution is added dropwise simultaneously with a reactionsolution to gradually change the pH of the solution from the acidic tothe alkaline range, thereby forming manganese ferrite nanoparticles inthe reaction field. Then, the solution and the manganese ferritenanopowder are separated, and the powder is dried to obtain a manganeseferrite powder having an average powder particle diameter of 1 nm ormore and less than 1000 nm (1 μm). The above method is an example of aninexpensive method because the steps are simple. In particular, all ofthe steps in the working examples of the present invention are carriedout at room temperature, and hence the burden of equipment costs andrunning costs in production steps is reduced due to the use ofproduction step that do not use a heat source. Although the method forproducing the manganese ferrite nanopowder used in the present inventionis of course not limited to the above-mentioned production method, theinitial liquid used in the above production method of the reaction fieldbefore the reaction starts (in the present application, this is alsoreferred to as the “reaction field solution”), the reaction solution,and the pH adjusting solution are now described in more detail below.

As the reaction field solution, an acidic solution is preferable. Inaddition to inorganic acids such as hydrochloric acid, nitric acid,sulfuric acid, and phosphoric acid, a solution obtained by dissolving ametal salt, a double salt thereof, a complex salt solution, and the likein a hydrophilic solvent such as water (e.g., an iron chloride solution,a manganese chloride solution, etc.), a solution of a hydrophilicsolvent such as an aqueous solution of an organic acid (e.g., aceticacid, oxalic acid, etc.), and combinations thereof, may be used. As thereaction field solution, preparing the reaction solution in advance inthe reaction field is effective for efficiently promoting the synthesisreaction of the manganese ferrite nanopowder. If the pH is less than −1,the material providing the reaction field is restricted, and avoidableimpurities may become mixed in the solution. Therefore, it is desirableto control the pH to between −1 or more and less than 7. To increase thereaction efficiency in the reaction field and minimize elution andprecipitation of unnecessary impurities, a particularly preferable pHrange is 0 or more and less than 7. As a pH range that provides a goodbalance between reaction efficiency and yield, the pH is more preferably1 or more and less than 6.5. Although hydrophilic solvents among organicsolvents and the like can be used as the solvent in the reaction field,it is preferable that water is contained so that the inorganic salt canbe sufficiently ionized.

The reaction solution may be a solution of an inorganic salt in water asa main component, such as a chloride such as iron chloride or manganesechloride, a nitrate such as iron nitrate, or a nitrite, a sulfate, aphosphate, or a fluoride containing an Fe component and/or a Mncomponent (optionally also containing an M component). In some cases, asolution mainly comprising a hydrophilic solvent, such as organic acidsalt in water may also be used as required. Also, a combination thereofmay be used. It is essential that reaction solution contain iron ionsand manganese ions. Regarding the iron ions, the reaction solution maycontain only divalent iron (Fe²⁺) ions, a mixture with trivalent iron(Fe³⁺) ions, or only trivalent iron ions. In the case of containing onlyFe³⁺ ions, it is necessary to contain metal ions of the M componentelement that are divalent or less. Representative examples of thevalence of the Mn ions are divalent, tetravalent, and heptavalent, butdivalent is best in terms of the homogeneity of the reaction in thereaction solution or reaction field solution.

Examples of the pH adjusting solution include an alkaline solution suchas sodium hydroxide, potassium hydroxide, sodium carbonate, sodiumhydrogencarbonate, and ammonium hydroxide, an acidic solution such ashydrochloric acid, and combinations thereof. It is also possible to usea pH buffer such as an acetic acid-sodium acetate mixed solution, or toadd a chelate compound or the like.

Although the oxidizing agent is not indispensable, it is an essentialcomponent when only Fe²⁺ ions are contained as Fe ions in the reactionfield solution or the reaction solution. Examples of the oxidizing agentinclude nitrites, nitrates, hydrogen peroxide, chlorates, perchloricacid, hypochlorous acid, bromates, organic peroxides, dissolved oxygenwater, and the like, and combinations thereof. Stirring in air or in anatmosphere having a controlled oxygen concentration is effective inmaintaining a situation in which dissolved oxygen acting as an oxidizingagent is continuously supplied to the manganese ferrite nanoparticlereaction field, and to control the reaction. In addition, bycontinuously or temporarily introducing an inert gas such as nitrogengas or argon gas by bubbling into the reaction field, for example, tolimit the oxidizing action of oxygen, the reaction can be stablycontrolled without inhibiting the effect of other oxidizing agents.

In a typical manganese ferrite nanopowder production method, formationof the manganese ferrite nanoparticles proceeds by the followingreaction mechanism. The nuclei of the manganese ferrite nanoparticlesare produced in the reaction solution directly or via an intermediateproduct such as green rust. The reaction solution contains Fe²⁺ ions,which are adsorbed on powder nuclei already formed or on OH groups onthe powder surface that have grown to a certain extent, therebyreleasing H⁺. Subsequently, when an oxidation reaction is performed byoxygen in the air, an oxidizing agent, an anode current (e⁺), or thelike, a part of the adsorbed Fe²⁺ ions is oxidized to Fe³⁺ ions. Whilethe Fe²⁺ ions, or the Fe²⁺ and the Mn²⁺ ions (or, Mn and M componentions), in the solution are again adsorbed on the already adsorbed metalions, H⁺ ions are released in conjunction with hydrolysis, whereby aferrite phase having a spinel structure is formed. Since OH groups arepresent on the surface of the ferrite phase, metal ions are againadsorbed and the same process is repeated to thereby grow into manganeseferrite nanoparticles.

Among these reaction mechanisms, to directly change from Fe²⁺ and Mn²⁺to the ferrite having a spinel structure, the reaction system may be,while adjusting the pH and the redox potential so as to cross the linedividing the Fe²⁺ ions and ferrite on the equilibrium curve in thepH-potential diagram of Fe, (slowly) shifted from the stable region ofFe²⁺ ions to the region where ferrite precipitates. Mn₂₊ is, except forspecial cases, a divalent state from the early stage of the reaction,and has almost no influence on redox potential change. In many cases,reactions due to a change in the redox potential of Fe (i.e., progressfrom the mixed solution to the ferrite solid phase) are described. Whenions of the M component element are contained and the oxidation numberof those ions changes and participates in the reaction, the sameargument can be made by using or predicting a pH-potential diagramcorresponding to the composition and the temperature. Therefore, it isdesirable to produce a ferrite phase while appropriately adjustingconditions such as the kind, concentration, and addition method of thepH adjusting agent and the oxidizing agent.

In most generally well-known ferrite nanopowder production methods, thereaction solution is prepared on the acidic side, the alkali solution isadded in one go to set the reaction field to a basic region, and fineparticles are instantaneously formed by coprecipitation. It may bethought that consideration is given such that differences in thesolubility product between the Fe component and the Mn component do notcause non-uniformity. Of course, the ferrite nanopowder may be preparedby such a method and very small nanoparticles can be prepared, and hencesuch a ferrite nanopowder can be used as the ferrite raw material forthe magnetic material of the present invention.

On the other hand, in the embodiment of the present invention, a step isdesigned such that, while dropping the reaction solution and supplyingthe raw materials for the manganese ferrite nanopowder production methodto the reaction field, the Mn component is steadily incorporated intothe Fe-ferrite structure by dropping the pH adjusting agent at the sametime to gradually change the pH from acidic to basic. According to thisstep, at the stage of producing the manganese ferrite nanopowder, the H⁺released when ferrite is produced by the above-mentioned mechanism isneutralized by the continuous introduction of the pH adjusting solutioninto the reaction field, and manganese ferrite particles are producedand grow one after another. Further, at the early stage of the reaction,there is a period in which green rust is produced and the reaction fieldbecomes green. However, it is important that the Mn component is mixedinto this green rust. When the green rust has finally been convertedinto ferrite, the Mn is incorporated into the lattice, and in thesubsequent reduction reaction, in the first phase and the second phase,the Mn is incorporated into the α-Fe phase having the bcc structure.

In addition to the above, other factors for controlling the reactioninclude stirring and reaction temperature.

Dispersion is very important to prevent the fine particles produced bythe manganese ferrite nanopowder synthesis reaction from agglomeratingand inhibiting a homogeneous reaction. To carry out such dispersion, anyknown method, or a combination thereof, may be used in accordance withthe purpose of controlling the reaction, such as a method in which thereaction is subjected to excitation while simultaneously dispersing byultrasonic waves, a method in which a dispersion solution is conveyedand circulated by a pump, a method of simply stirring by a stirringspring or a rotating drum, and a method of shaking or vibrating with anactuator or the like.

Generally, since the reaction in the manganese ferrite nanopowderproduction method used in the present invention is carried out in thepresence of water, as the reaction temperature, a temperature betweenthe freezing point and the boiling point of water under atmosphericpressure, namely, from 0° C. to 100° C., is selected.

In the present invention, a material produced from a method, e.g., asupercritical reaction method, for synthesizing manganese ferritenanopowder in a temperature range exceeding 100° C. by placing theentire system under high pressure may be, as long as a manganese ferritenanoparticles exhibiting the effects of the present invention can beformed, considered to be the magnetic material of the present invention.

As a method for exciting the reaction, in addition to theabove-mentioned temperature and ultrasonic waves, pressure and photoexcitation may also be effective.

Further, in the present invention, when applying a manganese ferritenanopowder production method using an aqueous solution containing Fe²⁺as the reaction solution (particularly when reacting the manganeseferrite nanopowder under conditions in which the Fe is mixed as adivalent ion), if the Mn content is less than 40 atom %, it is importantthat divalent ions of Fe are observed in the finally formed ferritenanopowder of the magnetic material of the present invention. The amountof the divalent ions is, in terms of the ratio of Fe²⁺/Fe³⁺, preferably0.001 or more. It is preferable to identify the divalent ions by usingan electron beam microanalyzer (EPMA). Specifically, the surface of themanganese ferrite nanoparticles is analyzed by the EPMA to obtain anX-ray spectrum of FeL_(α)-FeL_(β), the difference between the twomaterials is taken, and the amount of Fe²⁺ ions in the manganese ferritenanoparticles can be identified by comparing with the spectrum of astandard sample of an iron oxide containing Fe²⁺ (e.g., magnetite) andan iron oxide containing only Fe³⁺ (e.g., hematite or maghematite).

At this time, the EPMA measurement conditions are an accelerationvoltage of 7 kV, a measurement diameter of 50 μm, a beam current of 30nA, and a measurement time of 1 sec/step.

Examples of representative impurity phases of the manganese ferritenanopowder include oxides such as Mn-hematite, iron oxide hydroxidessuch as goethite, acagenite, lepidocrocite, feroxyhyte, ferrihydrite,and green rust, hydroxides such as potassium hydroxide and sodiumhydroxide, and the like. Among these, particularly when containing aferrihydrite phase and a Mn-hematite phase, since these form anα-(Fe,Mn) phase and other secondary phases after reduction, it is notalways necessary to remove them. These ferrihydrite and Mn-hematitephases are observed in SEM observation and the like as a sheet-likestructures having a thickness of several nm. However, since theparticles have a large area relative to their thickness, these phasesmay promote large improper grain growth in the reduction reactionprocess, and since they also contain many impurities other than the Fecomponent, the Mn component, and oxygen, it is desirable that the volumefraction of these phases is less than that of the manganese ferritenanopowder. In particular, when the atomic ratio of the Mn componentrelative to the Fe component is more than 0.33 and 0.5 or less, the Mnratio of the phases other than the manganese ferrite nanopowder centeredon ferrihydrite and Mn-hematite becomes larger than that of themanganese ferrite nanoparticles, and as a result, the disproportionationthat occurs during reduction becomes difficult to control. Therefore,careful attention needs to be given to the degree of agglomeration ofimpurity phases such as a ferrihydrite phase and a Mn-ferrite phase (inparticular, to prevent uneven distribution up to several microns). It isalso noted that, irrespective of the above, the ferrihydrite phase andMn-ferrite phase, which easily incorporate Mn, can be caused to coexistso as to prevent the above-mentioned inappropriate minor phases that donot contain Mn from precipitating by intentionally limiting the contentof these phases based on the whole magnetic material to a range from0.01% by volume to 33% by volume. When doing this, it is not necessaryto strictly maintain the control conditions during production of themanganese ferrite nanopowder, and hence the industrial benefits arelarge.

The average powder particle diameter of the manganese ferrite nanopowderused as a raw material of the present invention is preferably 1 nm ormore and less than 1 μm. It is more preferably 1 nm or more and 100 nmor less. If this average powder particle diameter is less than 1 nm, thereaction during reduction cannot be sufficiently controlled, resultingin poor reproducibility. If this average powder particle diameterexceeds 100 nm, the improper grain growth of the metal component reducedin the reduction step is substantial, and in the case of the softmagnetic material, the coercive force may increase, and hence theaverage powder particle diameter is preferably 100 nm or less. Further,if the average powder particle diameter is 1 μm or more, the α-Fe phaseseparates, Mn is not incorporated into this phase, and a magneticmaterial having poor in terms of the excellent electromagneticproperties and oxidation resistance provided by the present inventionmay be only obtained, and hence the average powder particle diameter ispreferably less than 1 μm.

When the manganese ferrite nanopowder used in the present invention isproduced mainly in an aqueous solution, moisture is removed bydecantation, centrifugation, filtration (in particular, suctionfiltration), membrane separation, distillation, vaporization, organicsolvent exchange, solution separation by magnetic field recovery of thepowder, or a combination thereof, and so on. The manganese ferritenanopowder is then vacuum dried at ordinary temperature or a hightemperature of 300° C. or less, or dried in air. The manganese ferritenanopowder may also be hot-air dried in air or dried by heat treating inan inert gas such as argon gas, helium gas, or nitrogen gas (in thepresent invention, the nitrogen gas may not be an inert gas depending onthe temperature range during heat treatment), or a reducing gas such ashydrogen gas, or a mixed thereof. Examples of a drying method thatremoves unnecessary components in the solution but does not use a heatsource at all include a method in which, after the centrifugation, thesupernatant is discarded, the manganese ferrite nanopowder is furtherdispersed in purified water, centrifugation is repeated, and finally thesolvent is exchanged with a hydrophilic organic solvent having a lowboiling point and a high vapor pressure, such as ethanol, and thenvacuum-dried under ordinary temperature.

(2) Reduction Step (in the present application, also referred to as“step (2)”)

This step is a step in which the manganese ferrite nanopowder producedby the above method is reduced to produce the magnetic material of thepresent invention.

Particularly preferably, the reduction step is carried out in a gasphase. Examples of the reducing atmosphere include hydrogen gas, anorganic compound gas, such as carbon monoxide gas, ammonia gas, andformic acid gas, a mixed gas of such an organic compound gas and aninert gas, such as argon gas and helium gas, a low-temperature hydrogenplasma, supercooled atomic hydrogen, and the like. Examples of methodsfor carrying out the reduction step include a method in which thesegases can be circulated in a horizontal or vertical tube furnace, arotary reaction furnace, a closed reaction furnace, or the like,refluxed, hermetically closed, and heated with a heater, and methods inwhich heating is carried out by infrared rays, microwaves, laser light,and the like. The reaction may also be carried out in a continuousmanner using a fluidized bed. Further, the reduction method such as themethod for reducing with solid C (carbon) or Ca, the method for mixingwith calcium chloride or the like and the method for reducing in aninert gas or a reducing gas, and as an industrial method, the method forreducing a Mn oxide with Al, may be used. However, as long as themagnetic material of the present invention is obtained, any method fallswithin the scope of the production method of the present invention.

However, a preferred method for the production method of the presentinvention is a method in which the reduction is carried out in hydrogengas or a mixed gas of hydrogen gas and an inert gas as the reducing gas.To produce the magnetic material of the present inventionphase-separated at the nano-scale, the reducing power is too strong byreducing with C or Ca, and it becomes very difficult to control thereaction for forming the soft magnetic material of the presentinvention. In addition, there are problems such as generation of toxicCO after reduction and mixing of calcium oxide, which must be removed bywashing with water. However, by reducing in hydrogen gas, the reductiontreatment can be carried out under consistently clean conditions.

However, conventionally, when Mn-ferrite is reduced with hydrogen gas,it was thought that (FeMn)O is formed when the temperature is increasedto 1000° C. (e.g., see Kiyoshi Terayama, Takayoshi Ishiguro, NetsuSokutei, Vol. 18, No. 3 (1991) pages 164-169 (hereinafter, referred toas Non Patent Document 5)). In addition, when C and hydrogen gas areused simultaneously, it is said that α-Fe and MnO are formed (e.g., seeT. Hashizume, K. Tarayama, T. Shimazaki, H. Itoh and Y. Okuno, J. ofThermal Analysis calorimetry, Vol. 69 (2002) pp. 1045-1050 (hereinafter,referred to as Non Patent Document 6)). However, the known literaturesuch as that mentioned above do not touch on the fact that α-(Fe,Mn) isformed when Mn-ferrite is reduced with hydrogen gas, and the applicantis unaware of any prior art literature disclosing that Mn²⁺ ions, whichare the main Mn ion in Mn-ferrite, are reduced to the valence of Mnmetal.

Further, even when considered thermodynamically, it is assumed from anEllingham diagram that Fe oxides are reduced in an H₂ gas flow, but Mnoxides are understood not to be easily reduced by H₂ gas. For example,at 1000° C., in the case of Fe, the ratio of H₂/H₂O when reduced frommagnetite to metallic iron is approximately 1, whereas for MnO it isaround 10⁵, from which it can be understood that even though the H₂ gasis flowing, there is little chance of the Mn oxide being reduced with H₂gas. Therefore, for a simple mixture or solid solution of an Fe oxideand a Mn oxide, it is usually considered that α-Fe and MnO or (FeMn)Oare formed by hydrogen reduction.

Therefore, the fact that the Mn²⁺ ion, which is the main Mn ion inMn-ferrite, is reduced to the valence of the Mn metal is believed tohave been first discovered by the present inventor. As to this fact, atthe present time, the following analogy can be made.

The Mn-ferrite of the present invention has a diameter of 1 nm or moreand less than 1000 nm (1 μm), Mn is atomically dispersed in a highlyactive nanopowder, and the affinity between Mn and Fe is high. As aresult, the Mn-ferrite is alloyed as α-(Fe,Mn) under a hydrogen gasflow. The reactivity of the nano-region particles is high, and contraryto thermodynamic expectations, often gives results beyond the commontechnical knowledge of the metallographic field. Conventionally, Mnoxides essentially cannot be reduced unless in the presence of Ca, Al,or the like, but according to the method of the present invention, apart of the Mn component is reduced to its metallic state and can bepresent as an alloy in the first phase or the α-(Fe,Mn) phase of thefirst phase and the second phase. At this time, the inventors infer thatthe coexistence of a small amount of an alkali metal such as Kinfluences the facilitation of the reaction.

The oxygen content in the material of the present invention is generallydetermined by an inert gas-melting method, but when the oxygen contentbefore reduction is known, the oxygen content in the material of thepresent invention can also be estimated from the weight differencebefore and after reduction. However, when there is simultaneously alarge amount of a halogen element, such as chlorine, whose content tendsto change before and after reduction, and an alkali element such as K orNa or a highly volatile component such as water or an organic componentcontained in the material, the content of each of these elements andcomponents should be individually identified. This is because the oxygencontent cannot be precisely estimated based only on the weight changebefore and after the reduction reaction.

Incidentally, among alkali metals derived from the raw materials, forexample, K begins to dissipate from the magnetic material at 450° C. dueto vaporization, and most of it is removed at 900° C. or above.Therefore, in the case of an alkali metal derived from the raw materialsfor which it is better to keep around in the early stage of thereduction reaction in order to utilize its catalytic action, butdepending on the application is preferably not present at the productstage, that alkali metal can be ultimately appropriately removed to anacceptable range by appropriately selecting the reduction conditions.The final content range of the alkali metal such as K that can be easilyremoved while having a positive effect on reduction is a lower limitvalue of 0.0001 atom % or more and an upper limit value of 5 atom % orless. This upper limit value can be further controlled to 1 atom % orless, and when most precisely controlled, to 0.01 atom %. Of course,based on the reduction conditions, it is also possible to reduce thealkali metal such as K further below the detection limit. Halogenelements such as Cl (chlorine) remaining in the manganese ferritenanopowder are mainly released outside the material system as hydrogenhalides such as HCl under the reducing atmosphere. The amount ofremaining Cl and the like starts to substantially decrease at areduction temperature of 450° C. or more, and although it depends on theMn and K contents and the content change thereof during the reductionstep, if a reduction temperature of approximately 700° C. or higher isselected, almost all of those halogen elements can be completely removedfrom inside the material.

The weight reduction before and after the reduction reaction of thepresent invention, which is mainly due to the O component beingconverted into H₂O and evaporating, is usually 0.1% by mass or more and80% by mass or less, although the amount depends on the Mn content, theM component content, the oxygen amount, the minor phase content, amountof impurities, amount of volatilized components such as water, thereducing reaction conditions such as the reducing gas species, and thelike.

Incidentally, as described in some of the Examples of the presentinvention, a local oxygen content may be determined based on aphotograph from an SEM and the like or by EDX, and a phase identified byXRD or the like may be specified on a microscopic observation image.This method is suitable for roughly estimating the oxygen content andits distribution in each phase of the first phase and the second phase.

Hereinafter, a method for producing the magnetic material of the presentinvention by a heat treatment in a reducing gas is described in detail.The heat treatment in a typical reduction step is carried out byincreasing the temperature of the material linearly or exponentiallyfrom room temperature to a constant temperature in a reducing gas flowat one or more temperature increasing rates, and then immediatelydecreasing the temperature linearly or exponentially to room temperatureusing one or more temperature decreasing rates, or maintaining thetemperature for a fixed period (=reduction time) when increasing ordecreasing the temperature during the temperature increasing/decreasingprocess or after the temperature has been increased (hereinafter,referred to as “constant temperature holding process). Unless statedotherwise, the reduction temperature of the present invention refers tothe highest temperature among the temperature at the time of switchingfrom the temperature increasing process to the temperature decreasingprocess and the temperature during the process of maintaining thetemperature for a fixed period.

When a method in which the manganese ferrite is reduced by hydrogen gasis selected as the production method of the soft magnetic material ofthe present invention, it is preferable to select a temperature range inwhich the reduction temperature is 400° C. or more and 1350° C. or less,although this depends on the Mn content. In general, this is becausewhen the reduction temperature is less than 400° C., the reduction rateis very slow, the reduction time is prolonged, and productivitydeteriorates. Further, when it is desired to reduce the reduction timeto one hour or less, it is preferable to set the lower limit of thereduction temperature to 500° C. or more.

When performing reduction at 1230° C. or more and 1538° C. or less,depending on the Mn content, the magnetic material being reduced maymelt. Therefore, generally if the Mn content is in the range of 0.01atom % or more and 33 atom % or less, the reduction treatment can becarried out by freely selecting the temperature range of approximately400° C. or more and 1350° C. or less. However, when the Mn contentexceeds 33 atom % and is up to 50 atom %, it is preferable to carry outthe reduction treatment by selecting a temperature of 400° C. or moreand 1230° C. or less.

A characteristic of the method for producing the magnetic material ofthe present invention is that since Mn is reduced to a metal stateaccording to the method of the present invention, performing thereduction reaction at the melting point or above, or at just below themelting temperature, may lead to coarsening of the microstructure, orthe Mn reacting with a reactor such as a ceramic container. From thisperspective, it is preferable not to set the reduction temperature to atemperature that is around or above the melting point. Depending on thecoexisting M component, it is generally desirable not to select atemperature higher than 1538° C. as the reduction temperature.

Further, since Mn has a high vapor pressure, during the reductionreaction, the amount of Mn vaporizing and being lost (here, sublimationis also referred to as “vaporization”) from the interior of thematerial, particularly the first phase, or additionally from the secondphase as well, increases at higher temperatures. Therefore, from thisperspective as well, when reducing in hydrogen at, for example, ordinarypressure at 800° C. or more, it is desirable to take measures such asadding Mn in excess relative to the target composition to be achieved.This is because such vaporization contributes to the disproportionationof the composition inside the material, which may be effective inreducing the coercive force.

From the above, the preferable reduction temperature range for themagnetic material of the present invention, which is a range in whichthe reduction time is short, the productivity is high, and the magneticmaterial does not melt, is 400° C. or more and 1350° C. or lessregardless of the Mn content. However, by permitting vaporization andloss of Mn from within the material and controlling the reductiontemperature to within a range of 800° C. or more and 1230° C. or less,it is possible to obtain the soft magnetic material of the presentinvention having an even smaller coercive force. Therefore, thistemperature range is particularly preferable in the present inventionfor producing a soft magnetic material having high magnetic properties.

When reduction is performed at the same temperature, the reductionreaction progresses as the reduction time increases. Therefore, thesaturation magnetization increases as the reduction time is longer, butfor coercive force, even if the reduction time is increased or thereduction temperature is increased, the coercive force does notnecessarily decrease. It is desirable to appropriately select thereduction time according to the desired magnetic properties.

When a method in which the manganese ferrite is reduced by hydrogen gasis selected as the production method of the semi-hard magnetic materialof the present invention, although this depends on the Mn content, it ispreferable to select the reduction temperature within a range of 400° C.or more and 1230° C. or less. This is because when the reductiontemperature is less than 400° C., the reduction rate may be very slow,the reduction time prolonged, and productivity deteriorate. On the otherhand, when the reduction temperature exceeds 1230° C., depending on theMn content, melting begins, so that the characteristics of thenanocrystals of the present invention may be inhibited and it may becomeimpossible to appropriately control the magnitude of the coercive force.A more preferable range of the reduction temperature is 450° C. or moreand 800° C. or less, and a particularly preferable range is 500° C. ormore and 700° C. or less.

As described above, when a method of reducing manganese ferrite withhydrogen gas is selected as the method for producing the magneticmaterial of the present invention, the preferable reduction temperaturerange is 400° C. or more and 1350° C. or less.

Since the magnetic material of the present invention contains Mn, thereduction rate is extremely slow compared with an intermediate ofFe-ferrite, for example, magnetite and maghemite (see Patent Document 1and Non Patent Document 3). For example, in the case of Fe-ferrite thathas an average powder particle diameter of 100 nm or less and that doesnot contain Mn, reduction to almost 100% by volume of α-Fe is carriedout just by reducing in hydrogen at 450° C. for 1 hour. Even underconditions of 425° C. for 4 hours, the Fe-ferrite is reduced to a levelat which it can hardly be observed even by X-ray diffraction.

On the other hand, when the reduction conditions are, for example, 450°C. for 1 hour, if the Mn content exceeds 20 atom %, the Mn-ferrite phasedisappears, and in XRD there is almost only the (Fe,Mn)O phase, and evenif an α-(Fe,Mn) phase is formed, the amount is slight. In addition, whenthe Mn content is 33 atom %, even when the temperature is increased to1200° C. and reduction is carried out for 1 hour, not only an α-(Fe,Mn)phase is present, the (Fe,Mn)O phase remains.

In the magnetic material of the present invention, due to such a slowreduction rate of the Mn-ferrite, reduction at a high temperature ispermitted. The nano-microstructure still containing Mn in the α-Fe phasedoes not become coarse, it can be formed into an aggregate ofmicrocrystalline structures comprising the first phase and the secondphase by a disproportionation reaction.

It is desirable that the magnetic material of the present invention isseparated into the first phase and the second phase at the nanoscale inthe reduction step during production of the magnetic material.Particularly in the case of the soft magnetic material of the presentinvention, it is necessary for the phases having the various Mn contentsand crystal structures to be separated by the disproportionationreaction, and for the orientation of those phases to be random or forthe phases to include concentration fluctuations in Mn content at thenanoscale. Further, it is also necessary for each of the crystal phasesto be ferromagnetically coupled.

The grains of the manganese ferrite nanoparticles grow as reductionprogresses. However, during that process, the crystal structure and theMn content of the first phase and the second phase, which are the formedcrystal phases, change in various ways depending on the reductiontemperature due to the Mn content of the original manganese ferritenanoparticles. In the temperature range of 400° C. or more and 912° C.or less, generally the Mn content of the first phase tends to decreaseas the temperature at which reduction to the metal phase occursincreases, and conversely, in the temperature range of 912° C. or moreand 1538° C. or less and the temperature range of room temperature ormore and less than 400° C., the Mn content of the first phase tends toincrease as the temperature is higher. However, when the reductionreaction is carried out at a temperature of 800° C. or more for a longtime, as described above, the Mn component is vaporized and released.Therefore, not only is the first phase influenced by the reduction time,but the Mn content as a whole is also influenced by the reduction time,and the Mn content tends to decrease as the reduction time becomeslonger.

Therefore, the composition of the crystal phases changes depending onthe rate of temperature increase during the increasing temperatureprocess and the temperature distribution in the reaction furnace.

In the initial stage of the reduction reaction, or when the reductiontreatment is carried out at a comparatively low temperature, forexample, when a Mn-ferrite nanopowder having a Mn content of 33 atom %and an average powder particle diameter of 100 nm or less is reduced at450° C. or less for 1 hour, the whole reduced powder first becomes an(Fe,Mn)O phase of 100 nm or less. Then, when the temperature isincreased to 500° C. or more, grain growth occurs by a plurality ofphases combining. Then, while reduction progresses, the phases begin toseparate due to the disproportionation reaction from the (Fe,Mn)Owustite phase mainly into α-(Fe,Mn) phases having a smaller Mn contentthan the original wustite phase and a wustite phase having a larger Mncomposition than the original wustite phase, and a powder having avariety of structures starts to appear as a results of the secondphases, such as the α-(Fe,Mn) phases having various compositions and the(Fe,Mn)O phases similarly having various compositions, piling on top ofeach other or many protrusions appearing on the surface of theparticles.

These various phase separation states are shown in FIG. 2. In FIG. 2(A),structures in which three phases having a Mn content in the range of30.0 atom % to 41.2 atom % are protruding like potato eyes fromparticles having a Mn content of 9.4 atom % can be seen. In FIG. 2(B),it can be seen that the Mn content of the three phases separated bygrain boundaries coming together at a grain boundary triple point isvery different in each phase, ranging from 11.9 atom % to 49.1 atom %.In the range of FIG. 2, it is identified from X-ray diffraction analysisand oxygen characteristic X-ray analysis that the phase having a Mncontent from 30.0 atom % to 49.1 atom % is a wustite phase.

As far as can be ascertained from SEM observation, the α-(Fe,Mn) phaseshaving a small Mn content undergo grain growth from tens of nm tohundreds of nm at times as the reduction temperature becomes higher. Inmany cases, (Fe,Mn)O phases with a size of several nm to several tens ofnm or less separate around there, forming structures bonded atinterfaces having a certain surface area while differing in phasecomposition and crystal structure. When decreasing the temperature, attemperatures of 400° C. or less, because the amount of Mn dissolved insolid solution in the α-(Fe,Mn) phases decreases, α-(Fe,Mn) phaseshaving a large Mn content begin to separate as the second phase from theα-(Fe,Mn) phases. Alternatively, γ-(Fe,Mn) phases, α-(Mn,Fe) phases, andβ-(Mn,Fe) phases having a large Mn solid solution range may also occur.For example, if the ferrite nanopowder of the present invention isreduced to a γ-(Fe,Mn) phase in a temperature range of 950° C. or morefor the whole composition region or a temperature range of 250° C. ormore in a given composition region, it can be inferred that the range inwhich the composition ratio of Fe and Mn is uniform is considerably wideas a result of the size of the nano region. When this powder material iscooled during the temperature decreasing process of the reductionreaction, the α-(Fe,Mn) phases having various Mn content undergo phaseseparation in accordance with the change in temperature. The magnitudeand the composition of that phase separation depend on the rate oftemperature decrease.

In these processes, although the α-(Fe,Mn) phase centered on the firstphase may grow even larger grains in some cases, the fact that numerouscrystal boundaries have occurred in the α-(Fe,Mn) phase whose grainshave been growing can be observed by SEM or TEM, thus showing thatinnumerable of microstructures having a crystal grain size of below 200nm are ubiquitous. An example of such a microstructure is shown in FIG.3. FIG. 3 is an SEM image of a magnetic material powder obtained byhydrogen reduction of a Mn-ferrite nanopowder at 1100° C., in which itcan be seen that the phases overall having an average Mn content of 0.7atom % are aggregates of fine crystal grains of 50 nm or less.

To show one example of a case in which, when the regions containing therespective crystal grains are observed by SEM-EDX, a ferrite nanopowderhaving a Mn content of 0.5 atom %, and in which there is a distributionin the Mn content, is reduced at 900° C. in a hydrogen gas flow, theα-(Fe,Mn) phase observed in a field of view of just 4 μm×3 μm exhibits awide Mn content distribution ranging from 0.3 atom % to 10 atom %. TheMn composition distribution in this α-(Fe,Mn) phase is also presumed tohave occurred by a disproportionation reaction of homologous phaseshaving the same crystal structure, and it is surmised that this reactionmainly occurs during the temperature decreasing process of the reductionstep. Therefore, it is very important to control the rate of temperatureincrease/decrease, including the temperature increasing process.Although the optimum conditions depend on the intended electromagneticcharacteristics and the starting composition, generally the rate oftemperature increase/decrease is appropriately selected between 0.1°C./min to 5000° C./min. When the Mn content is more than 20 atom %, itis preferable to control the rate of temperature increase/decrease to arate of 1° C./min to 500° C./min, since a soft magnetic material with alow coercive force can be prepared.

According to the equilibrium diagram of Fe—Mn, at around 500° C., up tonearly 3 atom % of the Mn can dissolve in solid solution in theα-(Fe,Mn) phase, but at ordinary temperature Mn hardly dissolves insolid solution in the α-Fe phase. The Mn content in the α-(Fe,Mn) phaseof the magnetic material of the present invention can be far beyond thissolid solution source, which is the equilibrium composition, butnaturally those will be non-equilibrium phases. If it were possible toperform the operation of decreasing the temperature from the reductiontemperature to ordinary temperature over infinite time (an infinitelysmall rate of temperature decrease), almost none of the Mn would coexistin α-Fe phase. Conversely, if it were possible to perform the operationof decreasing the temperature at an infinitely high rate from around thereduction temperature of 500° C. (an infinitely large rate oftemperature decrease), even if an α-(Fe,Mn) phase having a Mn content ofaround 3 atom % were present at the reduction temperature, α-(Fe,Mn)phases having various Mn contents due to a disproportionation reactiondo not separate from the α-(Fe,Mn) phase. Therefore, the soft magneticmaterial of the present invention cannot be obtained by the productionmethods for having any of the above limits. However, the magneticmaterial of the present invention has a microstructure that iscompletely different from existing materials in bulk, and does not havea composition distribution that follows the equilibrium state diagram atordinary temperature. However, near the reduction temperature,homogeneous phases in accordance with the equilibrium state spreadingacross the nano region in the magnetic material of the present inventionmay occur. In such a case, controlling the rate of temperatureincrease/decrease, including the temperature increasing process, may beimportant for the microstructure. From such a perspective, although theoptimum conditions depend on the intended electromagnetic properties andthe Mn content, it is desirable to appropriately select the rate oftemperature increase/decrease in the reduction step of the presentinvention generally between 0.1° C./min and 5000° C./min.

In addition, turning to the wustite phase, the XRD diffraction peakshifts to the lower angle side as the reduction temperature increases.This is because as the disproportionation reaction by reductionprogresses, the α-(Fe,Mn) phases having a small Mn content separate oneafter another from the wustite phase, so that Mn condenses into thewustite phase. As a result, the Mn content gradually increases, tellingus that it is approaching the manganosite phase.

As described above, when the ferrite nanopowder of the present inventionis reduced in hydrogen, a phase separation phenomenon due to thedisproportionation reaction very frequently occurs via the temperatureincreasing process, constant temperature maintenance process, andtemperature decreasing process, and during this period a wide variety ofphases having various compositions appear, whereby the magnetic materialof the present invention is formed.

In the present invention, the reason why appropriate grain growth occurswhile maintaining a nano-microstructure even in a high temperatureregion exceeding 800° C. is unknown. However, the raw material is amanganese ferrite nanopowder, and even if this is reduced by hydrogen toa metallic state like the first phase, as long as appropriate reductionconditions are selected, the original grain shape and compositiondistribution are not reflected whatsoever in the microstructure, thestructure has a uniform composition distribution, and there is noimproper grain growth like a coarsening of the crystal grain size. Sincethis grain growth occurs together with the reduction reaction, andconsidering that the volume reduction due to reduction is normally up to52% by volume, it can be easily inferred that disproportionationprogresses while leaving structures similar to intergrowths and skeletoncrystals. Further, it is also thought that, while the difference inreduction rates of the phases separated by disproportionation at theinitial stage of the reduction reaction is also involved, nanoscale veryfine disproportionated structures are ultimately formed as a whole dueto the phase separation caused by the disproportionation reaction duringthe temperature decreasing process mainly occurring in the α-(Fe,Mn)phase, causing nanoparticles and nanostructures to precipitate even fromthe high-temperature phases homogenized to a certain extent, which havea size in the nano region while maintaining their nano-microstructure.It is known that in the oxide phase containing Mn, such as theMn-ferrite phase, wustite phase, the reduction rate tends to be sloweras the Mn content is higher, and hence it is considered that oncedisproportionation occurs, the fact that the reduction reaction ratebecomes uneven within the material acts in a beneficial manner tomaintain the nanostructure.

The above series of observations is also supported by the fact that themagnetic material of the present invention should lose itscharacteristics if it melts.

(3) Gradual Oxidation Step (in the present application, also referred toas “step (3)”)

Since the magnetic material of the present invention after the reductionstep contains nano metal particles, there is a possibility that thematerial may spontaneously ignite and combust if directly exposed to theair. Therefore, although it is not an essential step, it is preferableto subject the magnetic material of the present invention to a gradualoxidation treatment immediately after the reduction reaction isfinished, as necessary.

The term “gradual oxidation” refers to suppressing rapid oxidation byoxidizing and passivating the surface of the reduced nano metalparticles (providing a surface oxide layer of wustite, Mn-ferrite,etc.). The gradual oxidation is carried out, for example, in a gascontaining an oxygen source, such as oxygen gas, in the vicinity ofordinary temperature to 500° C. or less, but in many cases a mixed gascontaining an inert gas with an oxygen partial pressure lower thanatmospheric pressure is used. If the temperature exceeds 500° C., itbecomes difficult to control and provide a thin oxide film of a few nmon the surface, no matter which low oxygen partial pressure gas is used.There is also a gradual oxidation method in which a vacuum is producedin a reactor, and then gradually released at ordinary temperature toincrease the oxygen concentration so that the reactor is not abruptlybrought into contact with the air.

In the present application, a step including an operation such as theabove is referred to as the “gradual oxidation step”. After this step,handling in the molding step, which is the next step, becomes verysimple.

Examples of a method for again removing the oxide film after this stepinclude a method in which the molding step is carried out under areducing atmosphere, such as hydrogen gas. However, since the surfaceoxidation reaction in the gradual oxidation step is not a completelyreversible reaction, it is impossible to remove all of the surface oxidefilm.

Of course, when the handling from the reduction step to the molding stepis carried out by an apparatus devised so that it can be operated in anoxygen-free state like a glove box, this gradual oxidation step isunnecessary.

Further, in the case of the magnetic material powder of the presentinvention, which has a large Mn content, a sufficiently high reductiontemperature and sufficiently long reduction time, and has undergonegrain growth, even if this magnetic material is exposed to the airwithout being subjected to this gradual oxidation step, stablepassivated films may be formed, and in such a case, a special gradualoxidation step is not required. In that case, exposing the magneticmaterial to the air can per se be regarded as a gradual oxidation step.

When oxidation resistance and magnetic stability are secured by gradualoxidation, ferromagnetic coupling may be broken by the oxide layer orthe layer of the passivated film, and hence it is preferable to performthe gradual oxidation after grain growth has occurred as much aspossible. Otherwise, as described above, it is preferable to not carryout the gradual oxidation step, and carry out the next molding step. Itis desirable to then continue the reduction step and the molding step bydeoxidation or a low oxygen process.

(4) Molding Step (in the present application, also referred to as “step(4)”)

The magnetic material of the present invention is used as a magneticmaterial (i.e., a solid magnetic material) in which the first phase andthe second phase are continuously bonded to each other directly or via ametal phase or an inorganic phase to form a massive state as a whole.The magnetic material powder of the present invention is used in variousapplications by solidifying the powder itself or by adding a metalbinder, another magnetic material, a resin, or the like and molding.When the magnetic material powder is in the state after step (2), orfurther after step (3), the first phase and the second phase may havealready been continuously bonded directly or via a metal phase or aninorganic phase. In this case, the magnetic material powder in thatstate functions as a solid magnetic material even without subjected tothe proper molding step.

As a method of solidifying only the magnetic material of the presentinvention, it is possible to use a method in which the magnetic materialpowder is placed in a mold, compacted in a cold state, and then used asit is, or the magnetic material powder may also be subjected to furthercold rolling, forging, shock wave compression molding and the like, andthen molded. In many cases, the method is carried out by sintering themagnetic material powder while heat treating it at a temperature of 50°C. or more. A method in which sintering is carried out withoutpressurization and just by heat treating is called pressurelesssintering method. The heat treatment atmosphere is preferably anon-oxidizing atmosphere, and it is desirable to perform the heattreatment in an inert gas, such as a rare gas like argon or helium ornitrogen gas, or in a reducing gas including hydrogen gas. The heattreatment can be carried out even in air if the temperature is 500° C.or less. Further, like pressureless sintering, the sintering may becarried out in a heat treatment atmosphere that is at ordinary pressure,or in a pressurized heat treatment atmosphere of 200 MPa or less, oreven in a vacuum.

Regarding the heat treatment temperature, in addition to ordinarytemperature molding carried out at less than 50° C., the heat treatmenttemperature is preferably 50° C. or more and 1400° C. or less forpressure molding and 400° C. or more and 1400° C. or less forpressureless sintering. At temperatures above 1230° C., the material maymelt, and hence it is preferable to carefully select the compositionrange. Therefore, a particularly preferable temperature range formolding is 50° C. or more and 1230° C. or less.

This heat treatment can also be carried out simultaneously with thepowder compacting. Further, the magnetic material of the presentinvention can be molded even by a pressure sintering method, such as hotpressing, HIP (hot isostatic pressing), electric current sintering, andSPS (spark plasma sintering). To make the pressurizing effect remarkablein the present invention, it is preferable that the pressurizing forcein the heating and sintering step is within the range of 0.0001 GPa ormore and 10 GPa or less. If the pressurizing force is less than 0.0001GPa, the effect of pressurization is poor and there is no change in theelectromagnetic properties from pressureless sintering. In such a case,pressure sintering is disadvantageous due to the resultant drop inproductivity. If the pressurizing force exceeds 10 GPa, the beneficiallimits of pressurizing are reached, and hence unnecessary pressurizingonly results in a drop in productivity.

In addition, strong pressurization imparts induced magnetic anisotropyto the magnetic material, and there is a possibility that thepermeability and coercive force deviate from the ranges in which theyare to be controlled. Therefore, the preferable range of thepressurizing force is 0.001 GPa or more and 2 GPa or less, and morepreferably 0.01 GPa or more and 1 GPa or less.

Among hot pressing methods, an ultra-high-pressure HP method, in which apowder compacted molded body is placed in a capsule that plasticallydeforms, and then hot pressed by heat treating while applying a strongpressure in one to three axis directions, is capable of inhibiting theentry of unwanted excess oxygen. This is because in such a method,unlike a hot pressing method in which the pressurized heat treatment isperformed in a die made of cemented carbide or carbon using a uniaxialcompressor, a pressure of 2 GPa or more, which is difficult even whenusing a tungsten carbide cemented carbide die, can be applied on thematerial without problems such as breaking the die, and the molding canbe carried out without contact with the air because the interior of thecapsule is hermetically sealed as a result of the plastic deformation bythe pressure.

Prior to molding, to adjust the powder particle diameter, coarsepulverization, fine pulverization, or classification can be carried outby using a known method.

Coarse pulverization is a step carried out before molding when thereduced powder is a massive object of several mm or more, or is a stepcarried out when again pulverizing after molding. Coarse pulverizationis carried out using a jaw crusher, a hammer, a stamp mill, a rotormill, a pin mill, a coffee mill, and the like.

Further, after coarse pulverization, in order to further adjust thedensity and molding properties at the time of molding, it is alsoeffective to adjust the particle diameter by using a sieve, a vibrationclassifier or sound classifier, a cyclone, and the like. Coarsepulverization and classification followed by annealing in an inert gasor hydrogen can eliminate structural defects and distortion, and in somecases may have an effect.

Fine pulverization is carried out when it is necessary to pulverize thereduced magnetic material powder or the molded magnetic material from asubmicron size to a size of several tens of μm.

Examples of the fine pulverization method include, in addition to themethods described above for coarse pulverization, using a dry or a wetfine pulverizing apparatus such as a rotary ball mill, a vibration ballmill, a planetary ball mill, a wet mill, a jet mill, a cutter mill, apin mill, and an automatic mortar.

A typical example of the method for producing the solid magneticmaterial of the present invention is to produce a manganese ferritenanopowder by step (1), reduce the manganese ferrite nanopowder by step(2), and then carry out step (3) followed by step (4), or performmolding only by step (4). A particularly preferable example of theproduction method is to prepare the manganese ferrite nanopowder by thewet method exemplified in step (1), then reduce the manganese ferritenanopowder by a method including hydrogen gas described in step (2),gradually oxidize the reduced manganese ferrite nanopowder to expose toa low oxygen partial pressure described in step (3) at ordinarytemperature, mold by the sintering method at ordinary pressure or underpressure described in step (4), in particular remove the oxygen on thepowder surface of the material in step (3), and then, as step (4), carryout molding in hydrogen to prevent any further oxygen from entering thematerial. The present solid magnetic material can be molded to athickness of 0.5 mm or more, and can be worked into an arbitrary shapeby cutting and/or plastic working.

When the magnetic material powder obtained by step (1)→step (2), or bystep (1)→step (2)→step (3), or by step (1)→step (2)→step (5) (describedlater), or by step (1)→step (2)→step (3)→step (5) (described later), orthe magnetic material powder obtained by re-pulverizing a magneticmaterial obtained by molding a magnetic material powder obtained by theabove steps by step (4), or the magnetic material powder obtained byannealing a magnetic material powder obtained by the above steps in step(5) (described later), is applied in a composite material with a resin,such as a high frequency magnetic sheet, the magnetic material powder ismolded by mixing with a thermosetting resin or a thermoplastic resin andthen compression molded, or is kneaded together with a thermoplasticresin and then injection molded, or is extrusion molded, roll molded,calendar molded or the like.

In the case of applying in an electromagnetic noise absorbing sheet, forexample, examples of the type of sheet shape include a batch type sheetobtained by compression molding, various rolled sheets obtained by rollmolding, calendar molding, and the like, and cut or molded sheets ofvarious sizes, such as A4 plate, having a thickness of 5 μm or more and10000 μm or less, a width of 5 mm or more and 5000 mm or less, and alength of 0.005 mm or more and 1000 mm or less.

(5) Annealing Step

The magnetic material of the present invention has a first phase and asecond phase, and typically one or both of those phases have a crystalgrain size in the nano region.

As long as the object of the present invention is not hindered, it maybe preferable to carry out annealing for various purposes, such as forcrystal distortions and defects that are produced in the various steps,stabilization of non-oxidized active phases, and the like. Theexpression “as long as the object of the present invention is nothindered” refers to the avoidance of situations in which thenanocrystals become more coarse due to, for example, improper graingrowth as a result of the annealing, or situations in which the magneticanisotropy near the crystal boundaries, which was required in order toadjust the permeability appropriately, is lost, thereby converselycausing an increase in the coercive force and inhibiting realization ofthe low permeability of the present invention.

For example, after the manganese ferrite nanopowder production step (1),to carry out stable reduction simultaneously with drying for the purposeof removing volatile components such as moisture content, a so-calledpreliminary heat treatment (annealing) in which fine particle componentsof about several nm are heat treated may be carried out for the purposesof inhibiting improper grain growth and removing lattice defects insubsequent steps. In this case, it is preferable to perform theannealing in air, in an inert gas, or in a vacuum at about 50° C. to500° C.

Further, the coercive force of the soft magnetic material of the presentinvention can be decreased by, after the reduction step (2), removingdistortions and defects in the crystal lattice and microcrystals causedby the decrease in the volume due to grain growth and reduction. Afterthis step, in applications in which the soft magnetic material of thepresent invention is used in powder form, for example, in applicationssuch as powder magnetic cores used by hardening a powder with a resin,ceramic, or the like, electromagnetic properties may be improved bycarrying out annealing under appropriate conditions after that step orafter a pulverization step or the like that is carried out after thisstep.

In addition, in the gradual oxidation step (3), annealing may be usefulfor removing distortions and defects caused by surface oxidation thatare present near the surface, interfaces, and boundaries.

Annealing after the molding step (4) is most effective. The annealingstep may be proactively carried out after preliminary molding,compression molding, hot pressing, and the like, or the subsequentcutting and/or plastic working to remove the distortions and defects inthe crystal lattices and microstructure caused by those steps. In theannealing step, there is expected to be a dramatic decrease in thedistortions, defects, and the like that have accumulated in the stepsprior to that. Furthermore, after the above-mentioned cutting and/orplastic working steps, the distortions in steps (1) to (4), steps (2) to(4), steps (3) and (4), or step (4) may be annealed, or the distortionsthat have accumulated in those steps may be annealed collectively.

The annealing atmosphere may be any one of a vacuum, a reduced pressure,an ordinary pressure, or a pressurized atmosphere of 200 MPa or less.The gas species to be used may be an inert gas, typified by a rare gassuch as argon, nitrogen gas, a reducing gas such as hydrogen gas, or anatmosphere containing an oxygen source such as air. The annealingtemperature may be from ordinary temperature to 1350° C., and in somecases the treatment may be carried out at a low temperature from aliquid nitrogen temperature to ordinary temperature. The apparatus usedin the annealing step may be the same as the apparatus used in thereduction step and the molding step, or it may be constructed bycombining known apparatuses.

EXAMPLES

The present invention will now be described in more detail by way ofexamples, but the present invention is in no way limited to theseexamples.

The methods for evaluating the present invention are as follows.

(I) Saturation Magnetization, Coercive Force, and Permeability

In the case of a magnetic powder, the powder was placed in a cylindricalcase made of polypropylene (inner diameter: 2.4 mm, powder layerthickness approximately 1.5 mm).

In the case of a molded powder, the molded body was molded on a diskhaving a diameter of 3 mm and a thickness of approximately 1 mm. Then,using a vibrating sample type magnetometer (VSM), a full loop of themagnetic curve in the region where the external magnetic field is −7.2MA/m or more and 7.2 MA/m or less was drawn, and the values of thesaturation magnetization (emu/g) and coercive force (A/m) at roomtemperature were obtained. The saturation magnetization was correctedwith a 5N Ni standard sample, and calculated based on the law ofapproach to saturation. The coercive force was corrected using aparamagnetic Pd standard sample to correct the magnetic field shift inthe low magnetic field region. In this measurement, if a smooth step orinflection point is not seen on the magnetic curve up to the zeromagnetic field after magnetization up to 7.2 MA/m, it is determined thatthere is no (i.e. “absent”) “inflection point on the ¼ major loop”. Thedirection of the measurement magnetic field is the axial direction inthe case of the magnetic powder and the radial direction in the case ofthe molded body.

When measuring a molded body, the saturation magnetization was convertedinto T (Tesla) units using the density of the molded body. The relativepermeability of the molded body was obtained by determining ademagnetizing factor based on a Ni standard sample having the same shapeas the measurement sample as mentioned above, and then roughlyestimating the relative permeability by using a magnetic curve having ademagnetizing field corrected based on the determined value.

(II) Oxidation Resistance

The saturation magnetization σ_(st) (emu/g) of a magnetic powder thathad been left in air at an ordinary temperature for a certain period t(days) was measured by the above method, compared with an initialsaturation magnetization σ_(s0) (emu/g), and the rate of decrease in thesaturation magnetization was evaluated based on the followingexpression.

Δσ_(s)(%)=100×(σ_(s0)−σ_(st))/σ_(s0)

The oxidation resistance performance can be determined as being higheras the absolute value of Δσ_(s) approaches zero. In the presentinvention, a magnetic powder having an absolute value of Δσ_(s) of 1% orless was evaluated as having good oxidation resistance for a period of tdays. In the present invention, t (days) is 60 or 120.

(III) Electric Resistivity

The molded body was measured by the van der Pauw method.

(IV) Fe Content, Mn Content, Oxygen Content, and α-(Fe,Mn) phase VolumeFraction

The Fe content and the Mn content in the powder and the bulk magneticmaterial were quantified by X-ray fluorescence elemental analysis. TheFe content and the Mn content in the first phase and the second phase ofthe magnetic material were quantified by EDX included in an FE-SEM orTEM based on an image observed by the FE-SEM or TEM. Further, the volumefraction of the α-(Fe,Mn) phases was quantified by image analysis bycombining a method using the above-mentioned FE-SEM together with theresults of the XRD method. Mainly to distinguish whether the observedphase is an α-(Fe,Mn) phase or an oxide phase, an oxygen characteristicX-ray surface distribution map using SEM-EDX was used. In addition, thevalidity of the value of the volume fraction of the α-(Fe,Mn) phases wasalso confirmed from the value of the saturation magnetization measuredin (I).

The oxygen content was determined based on a value obtained byquantifying by an inert gas melting method. Further, the oxygen contentof the magnetic material after the reduction step was also confirmedbased on the decrease in weight after reduction. In addition, imageanalysis by SEM-EDX was used for identification of each phase.

The K content was quantified by X-ray fluorescence elemental analysis.

(V) Average Powder Particle Diameter

The powder particle diameter was determined by observing the magneticpowder with a scanning electron microscope (SEM) or a transmissionelectron microscope (TEM). The powder particle diameter was determinedto one significant digit by selecting portions representing the wholematerial, and setting n to be a number of 100 or more.

When using together with a laser diffraction type particle sizedistribution analyzer, the volume-equivalent diameter distribution wasmeasured and evaluated in terms of a median diameter (μm) obtained fromthe distribution curve thereof. However, the value is employed only whenthe obtained median diameter is 500 nm or more and less than 1 mm. Itwas confirmed that such a value agrees to one significant digit with thepowder particle diameter roughly estimated by a method using amicroscope.

(VI) Average Crystal Grain Size

The magnetic material was observed with a scanning electron microscope(SEM), and the size of a portion surrounded by a crystal boundary wasobtained to one significant digit. The measurement area was determinedby selecting portions sufficiently representative of the whole, andsetting the number n to 100 or more. The crystal grain size wasdetermined by separately measuring the average value of the whole, andthe average value of only the first phase and the second phase,respectively. Further, the EDX device included with the transmissionelectron microscope (TEM) was used to investigate the size of theportions having a difference in Mn content and estimate the crystalgrain size at a fine scale. The number of measurement points of the Mncontent was set to 65536 points.

(VII) Crystallite Size

The crystallite size was determined by applying the Scherrer equation tothe line width of the (200) diffraction line of the bcc phases measuredby X-ray diffraction, and taking the dimensionless form factor to be0.9.

Example 1 and Comparative Example 1

An aqueous solution of MnCl₂.4H₂O (aqueous solution of manganese(II)chloride tetrahydrate) and an aqueous solution of FeCl₂.4H₂O (ferricchloride(II) tetrahydrate) were separately prepared, and then mixed toform a mixed aqueous solution of MnCl₂ and FeCl₂ adjusted to 50.3 mM,which was placed in a reactor as a reaction field solution. Next, a 660mM aqueous potassium hydroxide solution (pH adjusting solution) wasadded dropwise while vigorously stirring in air, and the pH of thesystem gradually shifted from the acidic side to the alkaline sidewithin a range of 4.64 or more and to 12.97 or less. At the same time, amixed aqueous solution (reaction solution) of FeCl₂ and MnCl₂ of 168 mMwas added dropwise and reacted for 15 minutes, then the addition of thepH adjusting solution and the reaction solution was stopped, and thestirring operation was further continued for 15 minutes. Next, the solidcomponent was precipitated by centrifugation, redispersed in purifiedwater and repeatedly subjected to centrifugation to adjust the pH of thesupernatant solution to 8.34. Finally, the precipitate was dispersed inethanol, and then subjected to centrifugation.

After that, vacuum drying was carried out at ordinary temperatureovernight to obtain a Mn-ferrite nanopowder having a(Fe_(0.672)Mn_(0.328))₄₃O₅₇ composition having an average powderparticle diameter of 20 nm. As a result of analyzing the nanopowder byX-ray diffraction, it was found that the cubic Mn-ferrite phase was themain phase and a slight amount of a rhombohedral Mn-hematite phase wascontained as an impurity phase. Further, an SEM image of this nanopowderis shown in FIG. 4. In the photograph, the spherical powder is aMn-ferrite nanopowder and the plate-like powder with a thickness ofseveral nm is the impurity phase. Therefore, this powder did not containan α-(Fe,Mn) phase, and was hence used as the powder of ComparativeExample 1. The particle diameter, magnetic properties, and the like ofthis powder are shown in Table 1.

The Mn-ferrite nanopowder was placed in a crucible made of aluminumtitanate, the temperature was increased at 10° C./min up to 300° C. in ahydrogen flow, then increased from 300° C. to 900° C. at 2° C./min, anda reduction treatment was carried out at 900° C. for 1 hour. After that,the temperature was lowered at a rate of 75° C./min to 400° C., and thencooled from 400° C. to room temperature over 40 minutes. Next, a gradualoxidation treatment was carried out at 20° C. in an argon atmospherehaving an oxygen partial pressure of 1% by volume for 1 hour to obtain amagnetic material having a composition ratio of manganese to iron ofFe_(67.0)Mn_(33.0). At this time, based on the whole magnetic materialincluding Mn, Fe, O, and K, the O content was 15 atom % and the Kcontent was zero. Further, the average powder particle diameter of theFe—Mn magnetic material was 30 μm. Analysis on this magnetic materialwas carried out by the following method, and this magnetic material wasused as Example 1.

As a result of observation of the obtained magnetic material by X-raydiffraction, it was confirmed that the α-(Fe,Mn) phase, which is a bccphase, is the main component. In addition, the presence of an (Mn,Fe)Owustite phase having a higher Mn content than this phase was alsoconfirmed. As a result, it was confirmed that the bcc α-(Fe,Mn) phasecorresponds to the first phase and the wustite phase corresponds to thesecond phase. In the wustite phase in this example, the peak position ofthe diffraction line lies between FeO wustite, which does not containMn, and MgO manganosite, and is expected to have a composition almostthe same as manganosite.

Further, the peak position of the (110) diffraction line of the firstphase observed by X-ray diffraction was, when measured with CoK α-rays,0.099 degrees to the lower angle side compared with an α-Fe powderprepared by the same method as described above except that the Mncomponent was not added and the reduction temperature was 450° C.Moreover, the line width was 0.29 degrees.

The Mn content of the magnetic material of this example was measured byusing XRD, which is generally said to be superior in terms of capturingthe macroscopic characteristics of the overall material. As a result,based on the fact that when Mn is in solid solution in the α-Fe phase,it is known that the peak position of the diffraction line shiftsgenerally to the lower angle side (this is not the case for a Mn contentrange from more than 3 atom % to less than 4 atom %), and based on thepeak position of the (110) diffraction line of the α-(Fe,Mn) phase, theline width thereof, and the literature values, the Mn content can beestimated to be about 30 atom % or less. From this, it could beconfirmed that the Mn content at the peak position of the diffractionline is about 5 atom %.

Therefore, in the present example, it was found that an α-(Fe,Mn) phasehaving a high Mn content was present as a second phase in addition tothe wustite phase.

The magnetic material powder was also observed by FE-SEM/EDX, which issuitable for finding the local Mn content of the magnetic material andthe presence and extent of disproportionation. As a result, as shown inFIG. 5, the content of Mn in each phase of the magnetic material (thenumerical values in the diagram are the Mn content in each phase,represented as the percentage value of the atomic ratio of Mn to the sumof Mn and Fe in each phase) was found to be distributed in a verydisproportionate manner in the range of from 8.1 atom % to 75.1 atom %.In addition, in FIG. 5, innumerable curved crystal boundaries curved atan interval in the order of tens of nanometers were also observed in aregion thought to be one α-(Fe,Mn) phase. Regarding the composition ofthe crystal phase deemed to be present in this region, which is a regionof a radius between 100 nm and 150 nm, the distribution of thiscomposition is a measurement result obtained by averaging thecomposition regarding its Mn content. It was found that in this averagedcomposition distribution of the α-(Fe,Mn) phase (in this example, the Mncontent of the wustite phase observed by EDX exceeds 50 atom %) greatlyvaried depending on the location, from 8.1 atom % to 28.9 atom %.Therefore, it is clear from these results that even in the α-(Fe,Mn)phase region, there are phases that can be distinguished based on Mncontent, for example, an α-(Fe,Mn) phase having a Mn content of 28.9atom %, which is twice or more and 10⁵ times or less the content of anα-(Fe,Mn) phase having a Mn content of 8.1 atom %, and which is 2 atom %or more and 100 atom % or less, namely, that regarding the α-(Fe,Mn)phases, a phase other than the first phase and that corresponds to thesecond phase is also present. It is also noted that in the region ofFIG. 5, no phase having a composition of about 5 atom % was observed byFE-SEM/EDX, but in view of the XRD measurement results, it is inferredthat such a phase would be observed if the measurement region was wider.

The average crystal grain size of the whole magnetic material was 90 nm.The crystal grain sizes of the first phase and the second phase were 100nm and 70 nm, respectively. In addition, observation of the crystalboundary vicinity at a magnification of 750,000 times confirmed that noheterogenous phases existed near these crystal boundaries.

According to such image analysis, X-ray diffraction, oxygen content, andthe like, the volume fraction of the bcc phase was estimated to be 57%by volume.

The saturation magnetization of this magnetic material was 137.7 emu/g,the coercive force was 8 A/m, and there was no inflection point on the ¼major loop.

Therefore, since the magnetic material of Example 1 has a coercive forceof 800 A/m or less, it was confirmed to be a soft magnetic material. Themeasurement results of the phases, composition, particle diameter, andmagnetic properties of this example are shown in Table 1.

Comparative Examples 2 to 4

Ferrite nanopowders were prepared in the same manner as in Example 1,except that the Mn component (aqueous solution of manganese chloride)was not added.

Fe metal powders were prepared in the same manner as in Example 1,except that the above ferrite nanopowders were used and the reducingconditions were 425° C. for 1 hour (Comparative Example 2), the sametemperature for 4 hours (Comparative Example 3), and 450° C. for 1 hour(Comparative Example 4).

The measurement results of the particle diameter and magnetic propertiesare shown in Table 1.

Note that in these metal powders, the magnetic properties dramaticallydeteriorate just by leaving in air at room temperature. Table 2 showsthe rate of change Δσ_(s) (%) of the saturation magnetization.

Comparative Example 5 and Examples 2 to 12

Magnetic materials of the present invention were produced in the samemanner as in Example 1, except that the reduction temperature and timewere set to the temperature (Comparative Example 5: 450° C., Examples 2to 12: range of 500° C. to 1200° C.) and time shown in Table 1, and arate v of temperature decrease (° C./min) until 400° C. was set to, whenthe reduction temperature is taken to be T (° C.), a velocityrepresented by the following relational expression.

[Expression 2]

v=0.1T−15  (2)

The measurement results of these phases, composition, particle diameter,and magnetic properties are shown in Table 1.

In all the examples, the presence of the first phase and the presence ofthe second phase were confirmed.

In Comparative Example 5, as a result of analysis by X-ray diffraction,it was found that an α-(Fe,Mn) phase and a Mn-ferrite phase were notpresent, and the structure was a single phase of a wustite phase. Thenumerical values shown in Table 1 are the analysis results of this(Fe,Mn)O wustite phase. The magnetic curve of Comparative Example 5 hasa shape as if a magnetic curve of a ferromagnetic material having asmall magnetization has been added to a paramagnetic material having alarge magnetic susceptibility, and the characteristics of a weakmagnetic material distinctly appear. It was found that the saturationmagnetization of the material of Comparative Example 5 is very low, witha value 1/7 or less that of Examples 2 to 12.

FIG. 6 is an XRD graph showing the results of Comparative Example 4(α-Fe powder), which is described later, Example 4 (reduction at 600°C.), Example 7 (reduction at 800° C.), and Example 9 (reduction at 1000°C.). It was confirmed that within the range shown in this graph, thediffraction peak of the (110) plane of the bcc phase shifted to a lowerangle as the reduction temperature increased. In particular, at 1000°C., two bcc phases were observed, and the presence of a bcc phase as thefirst phase and the presence of a “Mn-enriched bcc phase” as the secondphase were clearly distinguished. It is noted that the shoulderstructure on the high angle side is due to the Kα₂ rays. In general, theposition of the actual diffraction line shifts to the lower angle sidewhen the absorption of the Kα₂ rays is subtracted. In the presentinvention, based on a comparison with the magnetic material notcontaining the Mn component of Comparative Example 4, it was confirmedthat Mn was present in the bcc phases, and the average value of the Mncontent of the bcc phases was calculated to one significant digit basedon the magnitude of the low angle shift and the literature values,although it was assumed that at this magnitude of the low angle shift,the influence of the absorption of the Kα₂ rays is largely canceled out.It should be noted that the change in the diffraction position due tothe nanocrystals was similarly assumed as being canceled out by thesubtraction when calculating the magnitude of the low angle shift in theabove comparison, that measurement error from the stability of the XRDapparatus used in these examples was also continuously measured from asample for which comparison is to be made on the same day, and thediffraction peak position of a Si standard sample before and after themeasurement did not change, thereby ensuring the validity of the abovecomparison.

When two or more bcc phases are observed, the “Mn content of the bccphase roughly estimated based on the magnitude of the (110) low angleshift” listed in the table is calculated based on the diffraction angleat the maximum value of the diffraction peak observed at the lowestangle. Even for diffraction peaks other than the (110) diffraction peak,that is, “diffraction peak observed at the lowest angle”, the low angleshift occurs by containing Mn.

In addition, the K content relative to the whole magnetic materialincluding Mn, Fe, O, and K was 0 atom % (Examples 8 to 11) at areduction temperature of 850° C. or more. In the temperature range from500° C. to 800° C. (Examples 2 to 7 and Example 12), it was 0.1 atom %or less.

As can be seen by comparing Examples 3 and 12, it was found that the bccphase volume fraction and the saturation magnetization tended toincrease and the coercive force tended to decrease as the reduction timeincreased, even at the same reduction temperature.

The measurement results of the saturation magnetization and coerciveforce of Examples 1 to 11 and Comparative Example 5 are summarized inFIG. 7 with respect to the reduction temperature.

In addition, Table 2 shows the rate of change Δσ_(s) (%) of thesaturation magnetization in Examples 1, 4, and 9. The fact that Δσ_(s)is a negative value indicates that saturation magnetization is improvedafter leaving at an ordinary temperature as compared with immediatelyafter preparation of each magnetic powder. From the results of thistable, it was found that the oxidation resistance of the metal powdersof these examples is good at t=60 or 120.

Example 13 and Comparative Example 6

(Fe_(0.994)Mn_(0.006))₄₃O₅₇ Mn-ferrite nanopowders having an averagepowder particle diameter of 20 nm were prepared in the same manner,except that the composition ratio of MnCl₂.4H₂O and FeCl₂.4H₂O in themixed solution was changed. However, the pH in the reaction systemvaried from 4.56 to 13.14, and the pH at the time when the step ofwashing the solution remaining by centrifugation was completed was 7.53.FIG. 8 shows an SEM image of the nanopowder prepared in this way. As aresult of analyzing this material by X-ray diffraction, it was foundthat the cubic Mn-ferrite phase is the main phase and the rhombohedralMn-hematite phase is slightly contained as an impurity phase. Therefore,this nanopowder did not contain an α-(Fe,Mn) phase, and hence was usedas the powder of Comparative Example 6. The particle diameter, magneticproperties, and the like of this nanopowder are shown in Table 3.

The Mn-ferrite nanopowder was placed in a crucible made of aluminumtitanate, the temperature was increased at a rate of 10° C./min fromroom temperature up to 300° C. in a hydrogen flow, an annealingtreatment was carried out by providing a constant temperaturemaintenance process for stopping the increase in temperature for 15minutes at 300° C., then the temperature was again increased to 900° C.at a rate of 10° C./min, and a reduction treatment was carried out atthat temperature for 1 hour. After that, the temperature was lowered at75° C./min to 400° C., and then cooled from 400° C. to room temperatureover 40 minutes. Next, a gradual oxidation treatment was carried out at20° C. in an argon atmosphere having an oxygen partial pressure of 1% byvolume for 1 hour to obtain a magnetic material having a compositionratio of manganese to iron of Fe_(99.5)Mn_(0.5). FIG. 1 is an SEM imageof this Fe—Mn magnetic material powder. This magnetic material is a softmagnetic material as described below. Analysis on this magnetic materialwas carried out by the following method, and this magnetic material wasused as a powder in Example 13.

The O content based on the whole magnetic material including Mn, Fe, O,and K of the powder of Example 13 was 0.3 atom %, and the K content was0 atom %. Further, the average powder particle diameter of the Fe—Mnmagnetic material powder was 50 μm.

As a result of observing this magnetic material powder by X-raydiffraction, it was confirmed that an α-(Fe,Mn) phase, which is a bccphase, is the main component. In addition, an (Mn,Fe)O wustite phasehaving a higher Mn content than this phase was also confirmed to beslightly present. As a result, it was confirmed that the bcc α-(Fe,Mn)phase corresponds to the first phase and the wustite phase correspondsto the second phase.

Further, the position of the (110) diffraction line of the first phaseobserved by X-ray diffraction was, when measured with CoK α-rays,slightly to the lower angle side compared with an α-Fe powder preparedby the same method as described above except that the Mn component wasnot added and the reduction temperature was 450° C. Moreover, a gentle(110) diffraction line extending 0.478 degrees to the low angle regionand distinct from the diffraction line forming the maximum peak wasobserved.

From the X-ray diffraction results, it was found that the first phase,which is a bcc phase, contained up to nearly 15 atom % of Mn. Inaddition, as the second phase, it was found that α-(Fe,Mn) having a highMn content of 1 atom % or more, which is twice or more the Mn content(0.5 atom %) of the first phase, was present in the material of Example13.

The Mn content calculated from the maximum value of the magnitude of thelow angle shift at which the (110) diffraction line intensity is at amaximum value was about 1 atom %.

As a result of SEM-EDX analysis of the material of Example 13 by usingrepresentative locations, it was found that the Mn content wasdistributed in a very disproportionate manner from 0.03 atom % to 45.3atom %. Innumerable curved crystal boundaries (see FIG. 1) curved at aninterval in the order of tens of nanometers were also observed in aregion thought to be a bcc phase having a low Mn content. Regarding thecomposition of the crystal phase in this region, which is a region of aradius between 100 nm and 150 nm, the distribution of this compositionis a measurement result obtained by averaging the composition regardingits Mn content. It was found that this averaged composition distributionwas also very broad, from 0.03 atom % to 2.83 atom %, and greatly varieddepending on the location. The average value of the Mn contentdetermined by SEM-EDX was about 0.4 atom %. Therefore, it is clear fromthese results that even in the α-(Fe,Mn) phase region, there are phasesthat can be distinguished based on Mn content, for example, an α-(Fe,Mn)phase having a Mn content of 2.83 atom %, which is twice or more and 10⁵times or less the content of an α-(Fe,Mn) phase having a Mn content of0.03 atom %, namely, that regarding the α-(Fe,Mn) phases, a phase otherthan the first phase and that corresponds to the second phase is alsopresent.

The volume fraction of all the bcc phases, including these secondphases, was estimated to be about 99% by volume.

The average crystal grain size of the whole magnetic material of Example13 was about 200 nm. The crystal grain size of the first phase and thesecond phase was 200 nm and 100 nm, respectively. In addition,observation of the crystal boundary vicinity at a magnification of750,000 times confirmed that no heterogenous phases existed near thesecrystal boundaries.

The saturation magnetization of the magnetic material of Example 13 was219.1 emu/g, the coercive force was 200 A/m, and there was no inflectionpoint on the ¼ major loop. Since the magnetic material of this examplehas a coercive force of 800 A/m or less, it was confirmed to be a softmagnetic material. The saturation magnetization of this magneticmaterial exhibited a value that exceeded the mass magnetization (218emu/g) of α-Fe.

The measurement results of the phases, composition, particle diameter,and magnetic properties of these examples are shown in Table 3.

Examples 14 to 16

Magnetic materials of the present invention were produced in the samemanner as in Example 13, except that the Mn content was set to theamount shown in Table 3. In all the examples, the presence of the firstphase and the presence of the second phase were confirmed.

The measurement results of the phases, compositions, particle diameters,and magnetic properties of these examples are shown in Table 3.

It is also noted that the analysis of the Mn composition amount and thelike of the bcc phase was carried out in the same manner as in Example13.

In these examples, the K content relative to the whole magnetic materialincluding Mn, Fe, O, and K was 0 atom %.

Examples 17 and 18

Magnetic materials of the present invention were produced in the samemanner as in Example 9, except that the Mn content, the reductiontemperature, and the reduction time were set as shown in Table 4, andthe rate of temperature increase and rate of temperature decrease wereset as shown in Table 4. Further, in both the examples, the presence ofthe first phase and the presence of the second phase were confirmed. Themeasurement results of the phases, compositions, particle diameters, andmagnetic properties of these examples are shown in Table 4. Forcomparison, the results of Example 9 in Table 1 are also provided inTable 4.

The “fast” condition and the “slow” condition of the rate of temperatureincrease/decrease shown in Table 4 are as follows.

(Rate of Temperature Increase)

“Fast”: Temperature is increased at 10° C./min until the predeterminedreduction temperature.“Slow”: Temperature is increased at 10° C./min up to 300° C., and from300° C. to the predetermined reduction temperature is increased at 2°C./min.

(Rate of Temperature Decrease)

“Fast”: Until 400° C., temperature is decreased at a rate of temperaturedecrease of 85° C./min, and from 400° C. until ordinary temperature,temperature is decreased over 40 minutes.“Slow”: Until 300° C., temperature is decreased at 2° C./min, and from300° C. until ordinary temperature, temperature is decreased over 30minutes.

By comparing this example with Example 9, it can be seen that when theMn content is 31 atom % or more and 32 atom % or less, and the reductiontemperature is 1000° C., the coercive force decreases by setting therate of temperature increase and the rate of temperature decrease bothto fast. When either of the rate of temperature increase and the rate oftemperature decrease is set to slow, the semi-hard magnetic material ofthe present invention is obtained, but if the rate of temperatureincrease and the rate of temperature decrease are both set to fast, thecoercive force reaches the region of a soft magnetic material of 410A/m.

Example 19

The magnetic material powder of Example 9 was placed in a 3 ϕ cementedcarbide metal die made of tungsten carbide, and then electric currentsintering was carried out in a vacuum at 150° C. under 1.4 GPa.

Next, this electrically-sintered body was annealed in hydrogen at 1000°C. for 1 hour to prepare a solid magnetic material. The “slow” conditionwas selected as the rate of temperature increase and the “fast”condition was selected as the rate of temperature decrease. FIG. 9 is anSEM image of the surface of the solid magnetic material of this example.A large number of crystal boundaries were observed to be present in thesintered layer. FIG. 10 is an oxygen characteristic X-ray surfacedistribution map of the region of FIG. 9 obtained using SEM-EDX. In thewhite part, the oxygen content is high, and it can be seen that thispart is a wustite phase. In the other examples as well, a method that isbasically the same as this was employed when distinguishing whether thephase actually observed by SEM is an α-(Fe,Mn) phase or a wustite phase.

The measurement results of the phases, composition, particle diameter,magnetic properties, and electric resistivity of this solid magneticmaterial are shown in Table 5.

Further, this solid magnetic material is the soft magnetic material ofthe present invention, with a coercive force of 560 A/m, and asdescribed in the above Example 9, in a powder state, the solid magneticmaterial is a semi-hard magnetic material, with a coercive force of 1100A/m. The coercive force decreased as a result of the solidification ofthe powder due to ferromagnetic coupling caused by the sintering.

From analysis of the magnetic curve, it was found that the relativepermeability of the solid magnetic materials of Example 19 is in theorder of 10³ to 10⁴.

Comparative Example 7

The powder of Comparative Example 3 was placed in a cemented carbide diemade of tungsten carbide, and then electric current sintering wascarried out in a vacuum at 315° C. under 1.4 GPa. The electricresistivity of the obtained material was 1.8 μnm. The measurementresults of the particle diameter, magnetic properties, and electricresistivity of this solidified material are shown in Table 5.

According to Table 5, it can be seen that the solid magnetic material ofthe present invention described in Example 19 has an electricresistivity that is one order of magnitude higher than that of asolidified material not containing Mn, and furthermore, compared withthe 0.1 μΩm of pure iron and the 0.5 μΩm of an electromagnetic steelsheet, for example, which are existing materials, an electricresistivity higher by two orders of magnitude.

Example 20

A Mn-ferrite nanopowder having a (Fe_(0.999)Mn_(0.001))₄₃O₅₇ compositionhaving an average powder particle diameter of 20 nm was obtained in thesame manner as in Example 1. As a result of analyzing the nanopowder byX-ray diffraction, it was found that the cubic Mn-ferrite phase was themain phase and a slight amount of a rhombohedric Mn-hematite phase iscontained as an impurity phase.

The Mn-ferrite nanopowder was placed in a crucible made of aluminumtitanate, the temperature was increased at 10° C./min up to 300° C. in ahydrogen flow, then increased from 300° C. to 900° C. at 2° C./min, anda reduction treatment was carried out at 900° C. for 1 hour. After that,the temperature was lowered at a rate of 75° C./min to 400° C., and thencooled from 400° C. to room temperature over 40 minutes. Next, a gradualoxidation treatment was carried out at 20° C. in an argon atmospherehaving an oxygen partial pressure of 1% by volume for 1 hour to obtain amagnetic material having a composition ratio of manganese to iron ofFe_(99.9)Mn_(0.1). At this time, based on the whole magnetic materialincluding Mn, Fe, O, and K, the O content was 0.2 atom % and the Kcontent was zero. Further, the average powder particle diameter of theFe—Mn magnetic material was 100 μm. Analysis on this magnetic materialwas carried out by the following method.

As a result of observing this magnetic material powder by X-raydiffraction, it was confirmed that an α-(Fe,Mn) phase, which is a bccphase, is the main component. In addition, an (Mn,Fe)O wustite phasehaving a higher Mn content than this phase was also confirmed to beslightly present. As a result, it was confirmed that the bcc α-(Fe,Mn)phase corresponds to the first phase and the wustite phase correspondsto the second phase.

The Mn content calculated from the maximum value of the magnitude of thelow angle shift at which the (110) diffraction line intensity is amaximum value was about 1 atom %.

Further, the crystallite size calculated from the (200) diffraction linewidth was about 30 nm.

As a result of SEM-EDX analysis of the material of this example usingrepresentative locations, a result was obtained in which the Mn contentwas very disproportionized, being from 0.01 atom % to 0.27 atom %.Innumerable curved crystal boundaries curved at an interval in the orderof tens of nanometers were also observed in a region thought to be a bccphase having a low Mn content. Regarding the distribution, which is ameasurement result obtained by averaging the composition in the regionbetween the radius of 100 nm and 150 nm relative to the Mn content, ofthe crystal phases in this region, it was found that in this averagedcomposition distribution the crystal phases greatly varied depending onthe location, having a distribution from 0.01 atom % to 0.12 atom %. Theaverage value of the Mn content determined by SEM-EDX was about 0.04atom %. Therefore, it is clear from these results that even in theα-(Fe,Mn) phase region, there are phases that can be distinguished basedon Mn content, for example, an α-(Fe,Mn) phase having a 0.12 atom % Mncontent, which is twice or more and 10⁵ times or less the content of anα-(Fe,Mn) phase having a 0.01 atom % Mn content, namely, that regardingthe α-(Fe,Mn) phases, a phase other than the first phase and thatcorresponds to the second phase is also present.

The volume fraction of the all the bcc phases, including these secondphases, was estimated to be about 99.9% by volume.

The average crystal grain size of the whole magnetic material of thisexample was 300 nm. The crystal grain size of the first phase and thesecond phase was 300 nm and 200 nm, respectively. The reason why thesecrystal grain sizes are measured to be larger than the crystallite sizeof 30 nm is thought to be that because that SEM observation is used forthe measurement of the average crystal grain size in this example, anddue to the low resolution of the SEM, or due to the presence of crystalboundaries that cannot be observed by an SEM, the value is measured tobe larger than it actually is. However, even when SEM observation isused for the measurement of the average crystal grain size of thisexample, the size is 10 μm or less, and the average crystal grain sizeof this example is confirmed to be within the range of the magneticmaterial of the present invention.

The saturation magnetization of the magnetic material of this examplewas 219.1 emu/g, the coercive force was 8 A/m, and there was noinflection point on the ¼ major loop. Since the magnetic material ofthis example has a coercive force of 800 A/m or less, it was confirmedto be a soft magnetic material. The saturation magnetization of thismagnetic material exhibited a value that exceeded the mass magnetization(218 emu/g) of α-Fe.

FIG. 11 is a Mn characteristic X-ray surface distribution map based onTEM-EDX analysis of a cross section of a sample of the powder of thisexample cut to a thickness of about 100 nm. The black portion indicatesa low Mn content and the gray portion indicates a high Mn content. Inthis EDX analysis, the screen was divided into 256×256 20 nm-squarepixels, and a characteristic X-ray dose at the center of each pixel wasmeasured by narrowing the electron beam diameter to 0.2 nm. According tothis analysis, it can be generally said that even if the electron beamdiameter is narrowed to 0.2 nm, composition information on thesurroundings of about 1 nm is picked up.

FIG. 12 is a histogram representing the distribution of the Mn content.In FIG. 12, the Mn content of all the measurement points is divided intosix classifications, namely, 0 or more and less than 0.1 atom %(indicated as 0 to 0.1 on the horizontal axis in the graph), 0.1 or moreand less than 0.2 atom % (indicated as 0.1 to 0.2 on the horizontal axisin the graph), 0.2 or more and less than 0.3 atom % (indicated as 0.2 to0.3 on the horizontal axis in the graph), 0.3 or more and less than 0.4atom % (indicated as 0.3 to 0.4 on the horizontal axis in the graph),0.4 or more and less than 0.5 atom % (indicated as 0.4 to 0.5 on thehorizontal axis in the graph), and 0.5 atom % or more (indicated as 0.5+on the horizontal axis in the graph), and the number of pixels that eachclassification has is shown as a percentage (referred to as abundanceratio). The abundance ratio of pixels corresponding to the first phasehaving a Mn composition of 0 atom % or more and less than 0.1 atom % was63% of the total, and the abundance ratio of pixels having a Mn contentof at least twice or more that and having a Mn content of 0.2 atom % ormore corresponding to the second phase was 26% of the total. Therefore,it was found that the second phase makes up at least 26% or more of thetotal. In this way, it was found that the powder of Example 20 hadseparated into a first phase and a second phase based on adisproportionation reaction among the bcc phases.

In addition, in SEM, the crystal grain size is observed from 200 nm to300 nm or less, whereas in TEM-EDX analysis, the crystal grain size is 1nm or more and several tens of nanometers or less, which is close to thecrystallite size by XRD. Thus, a portion in which the Mn compositionfinely varies as described above may not be taken to be a crystalboundary in SEM, and hence the measured value of the crystal grain sizemay be higher by one order of magnitude or more. However, since thisparticle diameter range falls within the range of 1 nm or more and 10 μmor less for any measurement method, the material was confirmed to be themagnetic material of the present invention.

Further, although a microstructure of several tens of nanometers or lessseen in TEM-EDX analysis is ubiquitous, it was also found by electrondiffraction that there are portions in which the orientation of thecrystals is aligned within a deviation of about ±2 degrees in a range ofseveral hundred nm to several μm.

Example 21

FIG. 13 shows a Mn characteristic X-ray surface distribution mapobtained by TEM-EDX analysis carried out in the same manner as inExample 20 of a powder of Fe_(70.2)Mn_(29.8) prepared by the same methodas in Example 1. A black region indicated by reference numeral 21 in thediagram is a void of the sample. The pure white regions indicated byreference numeral 22 in the figure are wustite phases (which are eachvirtually identical to a MnO phase and thus hereinafter referred to as a“MnO phase”), and the gray and black mottled regions indicated byreference numeral 23 in the figure are metal phases (which are eachvirtually identical to a bcc phase and thus hereinafter referred to as a“bcc phase”). As can be seen from FIG. 13, in the bcc phase, it isconfirmed that there is fluctuation in the composition (concentration)of Mn due to the disproportionation reaction. As described later, thereis a portion (pixel) having a Mn content in the range from 0.02 atom %to 8.53 atom % in this bcc phase, and a first phase (Mn content of 0.02atom %), and a second phase having a Mn content that is twice or morethat of the first phase coexists in an adjacent manner to each other.

FIG. 14(A) is a diagram plotting the correlation between Mn content andFe content and FIG. 14(B) a diagram plotting the correlation between Mncontent and O content after, of the 256×256=65536 measurement pointsthat are the centers of the respectively 20 nm-square pixels obtained bycomposition analysis, the points whose Fe, Mn, and O (oxygen) contentwas each 0 atom % had been removed due to being considered as voids. InFIG. 14(A), the measurement points where the oxygen content is 0 atom %are points on the straight line where Fe content+Mn content=100.Further, in FIG. 14(B), those measurement points are a point on thehorizontal axis. In total, there were 12758 points having an oxygencontent of 0 atom %. Among those points, although there are also pointswhere the Mn content is very high, such as 57.1 atom % and 15.6 atom %,the Mn content is mostly a value between approximately 0 atom % and 8.53atom %. From the above results, it was found that Mn was dissolved insolid solution in the bcc phase, confirming that the results of the XRDpeak shift shown in the above examples were qualitatively supported.

The straight line indicated by the white dashed line in FIG. 14 connectsthe point where Mn is not present in the bcc phase with the point in theMnO phase (composition). This line is hereinafter referred to as the“bcc-MnO line”. Most of the nearly 60,000 composition analysis pointsare evenly present near this bcc-MnO line, so that regardless of whetherthose points are measurement points belonging to a region where the bccphase makes up the majority (region indicated by reference numeral 23 inFIG. 13) or a region where the MnO phase makes up the majority (regionindicated by reference numeral 22 in FIG. 13), it can be seen that mostof those points are a mixture of MnO and bcc (note that MnO is notpresent at measurement points where the oxygen content is 0 atom %). Itis noted that since the range being measured is in the range of 1 nmdiameter and the depth direction is about 100 nm, the bcc phase and MnOphase are present in an interwoven manner in this minute volume of 1 nm4) (diameter)×100 nm. Therefore, in the region of reference numeral 22in FIG. 13, although it is shown as a pure white phase due to the grayscale employed in the diagram, a minute bcc phase also exists in thisportion. Since the MnO phase is a second phase, it was found that thefirst phase and the second phase disproportionate and are separated atthe scale of several nm to several tens of nanometers.

It was confirmed that Mn dissolved in solid solution in a metallic statefrom the fact that most of the measurement points are biased towardshaving a larger Mn content and a smaller Fe content than the bcc-MnOline.

In SEM observation of this powder, the average crystal grain size wasabout 100 nm. On the other hand, similarly to in Example 20, withTEM-EDX analysis even for this example, the grain size of crystals withdifferent Mn compositions is from 1 nm to several tens of nanometers.Therefore, it was confirmed that, regardless of which measurement methodwas used, the powder was within the range of 1 nm or more and 10 μm orless, which is the range of the crystal grain size, and was within therange of the magnetic material of the present invention.

In FIG. 14, although there are points indicating compositions in whichoxygen is totally dominant, those are due to the electron beam skimmingclosely past in the vicinity of the surface of a void or a sample endportion.

In addition, the volume fraction of the bulk bcc phase of this examplepowder was 55% by volume, the O content was 14 atom %, the averagepowder particle diameter was 40 μm, the saturation magnetization was137.7 emu/g, the coercive force was 10 A/m, and there was no inflectionpoint on the ¼ major loop.

TABLE 1 Average FeMn bcc Phase Powder Reduction Reduction Powder VolumeParticle Temperature Time Mn Content Fraction O Content Diameter Example(° C.) (hours) (atom %) First Phase Second Phase (% by volume) (atom %)(nm) Example 1 900 1 33.0 α-(Fe, Mn) α-(Fe, Mn) phase, 57 15 30000 phasewustite phase Comparative — — 32.8 — — 0 57 20 Example 1 Comparative 4251 0.0 — — 29 41 100 Example 2 Comparative 425 4 0.0 — — 97 1.6 2000Example 3 Comparative 450 1 0.0 — — 99 0.7 2000 Example 4 Comparative450 1 33.4 — — 0 50 100 Example 5 Example 2 500 1 33.7 α-(Fe, Mn) α-(Fe,Mn) phase, 34 25 200 phase wustite phase Example 3 550 1 33.6 α-(Fe, Mn)α-(Fe, Mn) phase, 45 29 200 phase wustite phase Example 4 600 1 33.3α-(Fe, Mn) α-(Fe, Mn) phase, 51 17 300 phase wustite phase Example 5 6501 34.1 α-(Fe, Mn) α-(Fe, Mn) phase, 55 15 300 phase wustite phaseExample 6 700 1 33.5 α-(Fe, Mn) α-(Fe, Mn) phase, 58 15 300 phasewustite phase Example 7 800 1 32.0 α-(Fe, Mn) α-(Fe, Mn) phase, 59 14500 phase wustite phase Example 8 850 1 32.0 α-(Fe, Mn) α-(Fe, Mn)phase, 57 15 1000 phase wustite phase Example 9 1000 1 32.1 α-(Fe, Mn)α-(Fe, Mn) phase, 59 14 50000 phase wustite phase Example 10 1100 1 25.9α-(Fe, Mn) α-(Fe, Mn) phase, 61 13 70000 phase wustite phase Example 111200 1 17.8 α-(Fe, Mn) α-(Fe, Mn) phase, 65 11 100000 phase wustitephase Example 12 550 4 33.0 α-(Fe, Mn) α-(Fe, Mn) phase, 48 19 200 phasewustite phase Average First Phase Second Phase Presence/ bcc Phase MnCrystal Average Average Absence of Content (atom %) Grain Size ofCrystal Crystal Saturation Coercive Inflection Estimated From the WholeGrain Size Grain Size Magnetization Force Point on ¼ Magnitude of (110)Example (nm) (nm) (nm) (emu/g) (A/m) Major Loop Low Angle Shift Example1 90 100 70 137.7 8 Absent 5 Comparative 20 — — 29.7 2400 Absent —Example 1 Comparative 100 — — 85.9 12000 Absent — Example 2 Comparative2000 — — 214.6 3700 Absent — Example 3 Comparative 2000 — — 216.6 3200Absent — Example 4 Comparative 100 — — 12.5 14200 Absent 5 Example 5Example 2 200 200 200 91.3 8300 Absent 5 Example 3 200 300 200 113.84700 Absent 5 Example 4 300 300 200 124.1 2100 Absent 5 Example 5 200300 200 132.0 2200 Absent 5 Example 6 200 300 200 133.9 2400 Absent 5Example 7 100 100 100 136.9 880 Absent 7 Example 8 100 100 70 136.3 530Absent 5 Example 9 80 90 60 135.5 1100 Absent 7 Example 10 50 40 60138.9 760 Absent 9 Example 11 70 40 100 151.3 470 Absent 5 Example 12300 400 200 117.4 2800 Absent 4

TABLE 2 Average FeMn bcc Phase Powder Average Days Left at Powder VolumeParticle Crystal Saturation Room Mn Content Fraction Diameter Grain SizeMagnetization Temperature Δσ_(S) Example (atom %) (% by volume) (nm)(nm) (emu/g) (days) (%) Example 1 33.0 57 30000 90 137.7 120 0.1 Example4 33.3 51 300 300 124.1 60 −0.2 Example 9 32.1 59 50000 80 135.5 60 −0.3Comparative 0 29 100 100 85.9 60 5.4 Example 2 Comparative 0 97 20002000 214.6 60 19.0 Example 3 Comparative 0 99 2000 2000 216.6 120 27.2Example 4

TABLE 3 Average FeMn bcc Phase Powder Reduction Reduction Powder VolumeParticle Temperature Time Mn Content Fraction O Content Diameter Example(° C.) (hours) (atom %) First Phase Second Phase (% by volume) (atom %)(nm) Example 13 900 1 0.5 α-(Fe, Mn) α-(Fe, Mn) phase, 99 0.3 50000phase wustite phase Comparative — — 0.6 — — 0 57 20 Example 6 Example 14900 1 0.4 α-(Fe, Mn) α-(Fe, Mn) phase, 99 0.3 50000 phase wustite phaseExample 15 900 1 2.3 α-(Fe, Mn) α-(Fe, Mn) phase, 98 0.6 50000 phasewustite phase Example 16 900 1 5.4 α-(Fe, Mn) α-(Fe, Mn) phase, 92 2.340000 phase wustite phase Average First Phase Second Phase Presence/ bccPhase Mn bcc Phase Mn Crystal Average Average Absence of Content (atom%) Content (atom %) Grain Size of Crystal Crystal Saturation CoerciveInflection Estimated From Estimated From the Whole Grain Size Grain SizeMagnetization Force Point on ¼ Magnitude of (110) SEM-EDX Example (nm)(nm) (nm) (emu/g) (A/m) Major Loop Low Angle Shift Observation Example13 200 200 100 219.1 200 Absent 1 0.4 Comparative 20 — — 43.3 7200Absent — — Example 6 Example 14 300 300 200 220.1 340 Absent 2 0.3Example 15 300 300 200 214.6 440 Absent 2 2 Example 16 300 300 300 205.0540 Absent 5 5

TABLE 4 FeMn bcc Phase Reduction Reduction Rate of Rate of Powder VolumeTemperature Time Temperature Temperature Mn Content Fraction O ContentExample (° C.) (hours) Increase Decrease (atom %) First Phase SecondPhase (% by volume) (atom %) Example 9 1000 1 Slow Fast 32.1 α-(Fe, Mn)α-(Fe, Mn) phase, 59 14 phase wustite phase Example 17 1000 1 Fast Fast31.2 α-(Fe, Mn) α-(Fe, Mn) phase, 57 14 phase wustite phase Example 181000 1 Fast Slow 32.2 α-(Fe, Mn) α-(Fe, Mn) phase, 58 14 phase wustitephase Average Average First Phase Second Phase Presence/ bcc Phase MnPowder Crystal Average Average Absence of Content (atom %) ParticleGrain Size of Crystal Crystal Saturation Coercive Inflection EstimatedFrom Diameter the Whole Grain Size Grain Size Magnetization Force Pointon ¼ Magnitude of (110) Example (nm) (nm) (nm) (nm) (emu/g) (A/m) MajorLoop Low Angle Shift Example 9 50000 80 90 60 135.5 1100 Absent 7Example 17 40000 60 70 50 134.0 410 Absent 7 Example 18 50000 70 80 80136.0 860 Absent 7

TABLE 5 Average First Phase Second Phase FeMn bcc Phase Crystal AverageAverage Powder Volume Grain Size of Crystal Crystal Mn Content FractionO Content the Whole Grain Size Grain Size Example (atom %) First PhaseSecond Phase (% by volume) (atom %) (nm) (nm) (nm) Example 19 28.9α-(Fe, Mn) α-(Fe, Mn) phase, 58 14 100 100 100 phase wustite phaseComparative — — — 93 4.5 2000 — — Example 7 Presence/ bcc Phase MnAbsence of Content (atom %) Saturation Saturation Coercive InflectionElectric Estimated From Magnetization Magnetization Force Point on ¼Density Resistivity Magnitude of (110) Example (emu/g) (T) (A/m) MajorLoop (g/cm²) (μΩm) Low Angle Shift Example 19 135.2 1.08 560 Absent 6.3814 7 Comparative 208.7 1.84 3500 Absent 7.02 1.8 — Example 7

INDUSTRIAL APPLICABILITY

According to the magnetic material of the present invention, it ispossible to have a high magnetization and solve the problem of eddycurrent loss due to a high electric resistivity, which are contradictorycharacteristics for conventional magnetic materials, and yet haveexcellent electromagnetic properties that combine the merits of bothmetallic magnetic materials and oxide-based magnetic materials which donot require complicated steps such as lamination, as well as have stablemagnetic properties even in air.

The present invention relates to a soft magnetic material used intransformers, heads, inductors, reactors, cores (magnetic core), yokes,magnet switches, choke coils, noise filters, ballast, and the likemainly used for power devices, transformers, and informationcommunication related devices, as well as a motor or a linear motor fora rotary machines such as various actuators, voice coil motors,induction motors, reactance motors and the like, and in particular, asoft magnetic material used for a rotor, a stator, and the like, forautomotive drive motors exceeding 400 rpm, motors for industrialmachines such as power generators, machine tools, various generators,and various pumps, and motors for domestic electric appliances such asair conditioners, refrigerators, and vacuum cleaners.

The present invention also relates to a soft magnetic material used inantennas, microwave elements, magnetostrictive elements, magneticacoustic elements, and the like, as well as in sensors that employ amagnetic field, such as Hall elements, magnetic sensors, currentsensors, rotation sensors, and electronic compasses.

In addition, the present invention relates to a semi-hard magneticmaterial used in relays such as monostable and bistable electromagneticrelays, switches such as torque limiters, relay switches, and solenoidvalves, rotating machines such as hysteresis motors, hysteresis couplinghaving a brake functions and the like, sensors for detecting a magneticfield, a rotation speed, and the like, a bias of a magnetic tag, a spinvalve element, and the like, a magnetic recording medium or element suchas a tape recorder, a VTR, a hard disk, and the like.

Further, the present invention can also be used for high frequency softmagnetic and semi-hard magnetic materials for high frequencytransformers and reactors, as well as magnetic materials suppressingobstacles due to unnecessary electromagnetic interference, such aselectromagnetic noise absorbing materials, electromagnetic waveabsorbing materials, and magnetic shielding materials, materials forinductor elements such as noise removing inductors, RFID (RadioFrequency Identification) tag materials, noise filter materials, and thelike.

REFERENCE SIGN LIST

-   21 Void in sample-   22 Wustite phase (virtually identical to MnO phase)-   23 Metal phase (virtually identical to bcc phase)

1. A soft magnetic or semi-hard magnetic material, comprising a firstphase having crystals with a bcc structure containing Fe and Mn and asecond phase containing Mn, the second phase having a Mn content thatis, when a sum of the Fe and the Mn contained in the second phase istaken to be 100 atom %, larger than a Mn content when a sum of the Feand the Mn contained in the first phase is taken to be 100 atom %,wherein the second phase is a phase having crystals with a bcc structurecontaining Fe and Mn, and the Mn content is, when the sum of the Fe andthe Mn contained in the phase is taken to be 100 atom %, is an amount of2 times or more and 10⁵ times or less and/or 2 atom % or more and the100 atom % or less relative to the Mn content when the sum of the Fe andthe Mn contained in the first phase is taken to be 100 atom %.
 2. Themagnetic material according to claim 1, wherein the magnetic material issoft magnetic.
 3. The magnetic material according to claim 1, whereinthe first phase has a composition represented by a composition formulaFe_(100-x)Mn_(x) (where x is 0.001≤x≤33 in terms of atomic percentage).4. The magnetic material according to claim 1, wherein the first phasehas a composition represented by a composition formulaFe_(100-x)(Mn_(100-y)M_(y))_(x/100) (where x and y are 0.001≤x≤33 and0.001≤y≤50 in terms of atomic percentage, and M is one or more of Zr,Hf, Ti, V, Nb, Ta, Cr, Mo, W, Ni, Co, Cu, Zn, and Si).
 5. (canceled) 6.The magnetic material according to claim 1, wherein the second phasecomprises a Mn-ferrite phase.
 7. The magnetic material according toclaim 1, wherein the second phase comprises a wustite phase.
 8. Themagnetic material according to claim 1, wherein the phase havingcrystals with a bcc structure containing Fe and Mn has a volume fractionof 5% by volume or more based on the whole magnetic material.
 9. Themagnetic material according to claim 6, wherein the magnetic materialhas a composition in a range of, based on the composition of the wholemagnetic material, 20 atom % or more and 99.998 atom % or less of Fe,0.001 atom % or more and 50 atom % or less of Mn, and 0.001 atom % ormore and 55 atom % or less of
 0. 10. The magnetic material according toclaim 1, wherein an average crystal grain size of the first phase, thesecond phase, or the whole magnetic material is 1 nm or more and 10 μmor less.
 11. The magnetic material according to claim 1, wherein atleast the first phase has a bcc phase having a composition representedby a composition formula Fe_(100-x)Mn_(x) (where x is 0.001≤x≤1 in termsof atomic percentage), and the bcc phase has a crystallite size of 1 nmor more and less than 100 nm.
 12. The magnetic material according toclaim 1, wherein the magnetic material is in a powder form, and when themagnetic material is soft magnetic, the magnetic material has an averagepowder particle diameter of 10 nm or more and 5 mm or less, and when themagnetic material is semi-hard magnetic, the magnetic material has anaverage powder particle diameter of 10 nm or more and 10 μm or less. 13.The magnetic material according to claim 1, wherein at least the firstphase and the second phase are ferromagnetically coupled with adjacentphases.
 14. The magnetic material according to claim 1, wherein thefirst phase and the second phase are continuously bonded to each otherdirectly or via a metal phase or an inorganic phase to form a massivestate as the whole magnetic material.
 15. A method for producing themagnetic material according to claim 12, comprising reducing a manganeseferrite powder having an average powder particle diameter of 1 nm ormore and less than 1 μm in a reducing gas containing hydrogen gas at areduction temperature of 400° C. or more and 1350° C. or less.
 16. Amethod for producing the magnetic material according to claim 1,comprising reducing a manganese ferrite powder having an average powderparticle diameter of 1 nm or more and less than 1 μm in a reducing gascontaining hydrogen gas, and forming the first phase and the secondphase by a disproportionation reaction.
 17. A method for producing themagnetic material according to claim 14, comprising reducing a manganeseferrite powder having an average powder particle diameter of 1 nm ormore and less than 1 μm in a reducing gas containing hydrogen gas at areduction temperature of 400° C. or more and 1350° C. or less to producea magnetic material, and sintering the magnetic material.
 18. A methodfor producing a soft magnetic or semi-hard magnetic material, comprisingproducing a magnetic material according to claim 15, and performingannealing at least once after the reduction step.
 19. A method forproducing a soft magnetic or semi-hard magnetic material, comprisingproducing a magnetic material according to claim 16, and performingannealing at least once after the reduction step.
 20. A method forproducing a soft magnetic or semi-hard magnetic material, comprisingproducing a magnetic material according to claim 17, and performingannealing at least once after the reduction step.
 21. A method forproducing the magnetic material according to claim 14, comprisingreducing a manganese ferrite powder having an average powder particlediameter of 1 nm or more and less than 1 urn in a reducing gascontaining hydrogen gas to produce a magnetic material, forming thefirst phase and the second phase by a disproportionation reaction, andsintering the magnetic material.